Compound semi-conductors and controlled doping thereof

ABSTRACT

A method of controlling the amount of impurity incorporation in a crystal grown by a chemical vapor deposition process. Conducted in a growth chamber, the method includes the controlling of the concentration of the crystal growing components in the growth chamber to affect the demand of particular growth sites within the growing crystal thereby controlling impurity incorporation into the growth sites.

The invention relates to the controlled growth of high-qualitysemiconductor device crystal films, and more particularly, to a methodand system for producing high-quality silicon carbide semiconductorswhich are highly reproducible. This invention was made with governmentsupport under a government contract awarded by NASA. The government hascertain rights in the invention.

This application is a continuation in part of U.S. patent applicationSer. No. 276,599 filed on Jul. 18, 1994 and issued as U.S. Pat. No.5,463,978 issued on Nov. 7, 1995, which is a continuation of U.S.application Ser. No. 008,650 filed Jan. 25, 1993, now abandoned.

INCORPORATION BY REFERENCE

U.S. Pat. No. 5,363,800 entitled "Process For The Controlled Growth OfSingle-Crystal Films Of Silicon Carbide Polytypes On Silicon CarbideWafers" issued Nov. 15, 1994 and U.S. Pat. No. 5,248,385 entitled"Process For The Homoepitaxial Growth Of Single-Crystal Silicon CarbideFilms On Silicon Carbide Wafers" issued Sep. 28, 1993 which bothillustrate the pretreatment of a silicon carbide substrate areincorporated herein by reference. U.S. Pat. No. 5,463,978 entitled"Compound Semicondctor And Controlled Doping Thereof" issued Nov. 7,1995 is also incorporated herein to illustrate the technique of sitecompetition epitaxy for crystals having two or more components.

BACKGROUND OF THE INVENTION

The invention is particularly applicable to production of siliconcarbide crystals (herein used to include crystal films) and it will bediscussed with particular reference thereto; however, the invention hasmuch broader applications and can be used for other crystals grown bythe chemical vapor deposition process.

Semiconductor devices are used in a wide variety of electronicapplications. Semiconductor devices include diodes, transistors,integrated circuits, light-emitting diodes and charge-coupled devices.Various semiconductor devices using silicon or compound semiconductorssuch as gallium arsenide (Gas) and gallium phosphide (GAP) are commonlyused. In order to fabricate semiconductor devices, it is necessary to beable to grow high-quality, low-defect-density single-crystal films withcontrolled impurity incorporation (with respect to both netconcentration and concentration profiles) while possessing good surfacemorphology. In recent years, there has been an increasing interest inresearch of the silicon carbide semiconductors for use in hightemperature, high powered and/or high radiation operating conditionsunder which silicon and conventional III-V semiconductors cannotadequately function.

Silicon carbide has been classified as a compound semiconductor withpotentially superior semiconductor properties for use in applicationsinvolving high temperature, high power, high radiation and/or highfrequency. Silicon carbide has a number of characteristics that make ithighly advantageous for various uses. Such advantages include a wideenergy gap of approximately 2.2 to 3.3 electron volts, a high thermalconductivity, a low dielectric constant, a high saturated electron driftvelocity, a high breakdown electric field, a low minority carrierlifetime, and a high disassociation temperature. Furthermore, siliconcarbide is thermally, chemically and mechanically stables and has agreat resistance to radiation damage. In addition, a variety of opticaldevices, such as light-emitting diodes (LEDs) can be fabricated fromsilicon carbide and operated at temperatures exceeding 600° C. Despitethese many advantages and capabilities of silicon carbide semiconductordevices, large scale commercialization of silicon carbide devices hasbeen slow because of the lack of control over the crystal quality,growth reproducibility, and controlled dopant incorporation into thesilicon carbide crystals.

Several properties of SiC contribute to this lack of control. First, itdoes not melt at reasonable pressures and it sublimes at temperaturesabove 1800° C. Second, it grows in many different crystal structures,called polytypes. Third, post-growth doping attempts (i.e. diffusioninto the crystal from a gas phase species as is used in the siliconindustry) are not effective in SiC crystals. Other known post-growthdoping techniques (i.e. ion implantation) typically result inconsiderable crystal damage. Attempts to remove this crystal damage, andtherefore improve device performance by post-annealing, commonly resultin severe dopant profile redistribution.

Since molten-SiC growth techniques cannot be applied to SiC, twotechniques have been developed to grow silicon carbide crystals. Thefirst technique is known as chemical vapor deposition (CVD) in whichreacting gases are introduced into a crystal chamber to form siliconcarbide crystals upon an appropriate heated substrate. A secondtechnique for growing bulk silicon carbide crystals is generallyreferred to as the sublimation technique or Lely process. In thesublimation technique, some type of solid silicon carbide material otherthan the desired single crystal in a particular polytype is used as astarting material and heated until the solid silicon carbide sublimes.The vaporized material is then condensed to produce the desiredcrystals. Although a large number of crystals can be obtained by eitherthe sublimation method or the epitaxial growth method (CVD), it isdifficult to prepare large single crystals of silicon carbide and tocontrol with high accuracy the size, shape, polytype and doping of thesilicon carbide crystals.

Silicon carbide crystals exist in hexagonal, rhombohedral and cubiccrystal structures. Generally, the cubic structure, with the zincblendestructure, is referred to as the β-SiC or 3C--SiC whereas numerouspolytypes of the hexagonal and rhombohedral forms are collectivelyreferred to as α-SiC. The most common α-SiC polytype is 6H--SiC. Each ofthe various silicon carbide polytypes have unique electrical and opticalproperties which give them advantages over the other polytypes inparticular applications. For example, the 6H--SiC polytype has a bandgapof about 2.9 electron volts and a hexagonal structure, wherein the3C--SiC polytype has a lower bandgap of about 2.2 electron volts and hasa higher symmetry structure than the 6H--SiC polytype. These propertydifferences lead to advantages for the 6H--SiC polytype in someapplications such as a wider bandgap resulting in blue light-emittingdiodes and operation at higher temperatures. On the other hand, the3C--SiC polytype has a higher electron mobility leading to a higherfrequency of operation.

SiC polytypes are formed by the stacking of double layers of Si and Catoms. Each double layer may be situated in one of three positions. Thesequence of stacking which determines the particular polytype. Thestacking direction is called the crystal c-axis which is perpendicularto the basal plane. For 6H--SiC polytypes, the (0001) plane (theSi-face) or (0001) plane (the C-face) is known as the basal plane andfor 3C--SiC, the plane (111) is equivalent to the (0001) basal plane.

Many advances have been made in the growing of higher quality 6H--SiCand 3C--SiC crystals having fewer dislocations, stacking faults,microtwins, double positioning boundaries (DPBs), threading dislocationsand anti-phase boundaries (APBs).

U.S. Pat. No. 5,037,502 (Suzuki) discloses a process for producing asingle-crystal silicon carbide by growing in a CVD chamber a singlecrystal film of α-silicon carbide on a single crystal film of β-siliconcarbide. Suzuki discloses that the CVD reaction chamber should bemaintained between 1400-1900° C. and preferably 1500°-1700° C. while theβ-silicon carbide crystal is growing. Silane SiH₄) and propane (C₃ H₈)are the source gases fed into the growth chamber at Si/C atomic numberratio of 0.01-10 and preferably 0.5-5. The Si/C ratio is set at aconstant rate during crystal growth. The setting of the Si/C ratio at asingle optimum point maximizes crystal growth while minimizing crystaldefects to grow the β-silicon carbide crystal. The α-silicon carbidecrystal is grown on top of the β-silicon carbide crystal by maintaininga similar Si/C ratio and reducing the growth chamber temperature to800-1200° C. and preferably 1000-1100° C. Suzuki discloses thatβ-silicon carbide crystals must grow at temperatures between 1200-1400°C. and preferably 1300-1350° C. By lowering the chamber temperaturebelow 1200° C., only α-silicon carbide crystals can grow. Suzuki onlyteaches that temperature affects the type of silicon carbide crystalwhich can be grown. Suzuki does not disclose (1) any method of dopingthe β or α-silicon carbide crystal, (2) any type of pretreatment processsuch as preparing the β-silicon carbide crystal prior to growing theα-silicon carbide crystal, (3) whether the tilting of the β-siliconcarbide substrate will affect the growth of the α-silicon carbidecrystal, and (4) whether a particular side of the β-silicon carbidesubstrate grows better α-silicon carbide crystals.

With all the advances in growing 3C--SiC and 6H--SiC crystals, thecontrolled purity and doping of such crystals are limited. Furthermore,controlled and reproducible degenerate doping or very low doping has ofyet been unachievable. Dopant incorporation in a grown crystal is knownas intentional impurity incorporation and, contaminant incorporationinto the crystal is known as unintentional impurity incorporation (i.e.contamination). During the sublimation or CVD crystal growing process,various compounds and/or elements are intentionally and unintentionallyincorporated into the crystal. Various techniques of limited successhave been used to exclude contaminants from a crystal to produce highlypure crystals.

One technique for doping silicon carbide crystals is disclosed in anarticle by H. J. Kim and R. F. Davis entitled Theoretical and EmpiricalStudies of Impurity Incorporation into β-SiC Thin Films During EpitaxialGrowth (1986). In the Kim reference, a study relating to the dopantincorporation of β-SiC crystals are disclosed. The Kim reference teachesthat the dopant incorporation in a β-SiC crystal increases as the dopantconcentration increases within the SiC growth reactor. Kim also teachesthat some dopants are incorporated in concentrations which are higherthan predicted from theory. Kim's observation of an unexpectedly highconcentration of incorporation, as determined by comparisons of SIMS andelectric probe results are hypothesized to be due to threepossibilities, namely, (I) line or point defects and/or trapping ofimpurities at dislocations and stacking faults known to be in thematerial, (ii) dopant-Si and/or dopant-C interaction, and (iii) locationin non-electrically active interstitial sites. The dopant-Si and/ordopant-C interaction that Kim refers to relates to a portion of each ofthe dopants either 1) not ionized or located on non-electrically activesites or 2) form complexes with Si or C. Kim's phrase "forms complexeswith Si or C" teaches that the dopant atom can combine with available Siand/or C to form distinctly different compounds possessing anon-electrical character within the SiC crystal, thereby creating thediscrepancy between the electrically measured dopant concentrationversus the atomically probed (i.e. SIMS) elemental dopant concentration.These proposed complexes are inclusions of compounds, such as thecompound silicon nitride (Si₃ N₄). The high melting point material beingreferred to in the Kim reference is silicon nitride (Si₃ N₄), which hasa melting of 1900° C. With regard to the location of non-electricallyactive interstitial sites, Kim suggests that the dopant atoms could alsobe located in other non-electrically active sites. Specifically, Kimmakes reference to "a non-electrically active interstitial" position inthe crystal lattice. Kim teaches that some atoms (e.g., Al and P),follow ideal incorporation behavior while others (i.e. N, B) apparentlydo not because these other dopant atoms are incorporated into the grownSiC crystal at a much higher concentration than predicted from theory.In summary, Kim teaches that the excess dopant incorporation in thedisclosed β-SiC crystal is primarily believed, as theorized by Kim, toresult from (1) accumulation of non-electrically active dopant atoms atdefect sites such as dislocations and stacking faults, (2)non-electrically active high melting point compounds formed frominteractions of dopant atoms with the SiC and/or C or (3) accumulationof dopant atoms at non-electrically active interstitial sites within theSiC lattice. These three possible explanations are well known to the artof SiC growth since the high density of defects in β-SiC (grown onsilicon substrates) is one of the main reasons it has been largelyabandoned for growth on the superior SiC substrates. Kim does notdisclose any techniques for controlling the rate of incorporation ofdopant atoms in a SiC crystal by varying the Si/C source gas ratio. Thestudy in Kim was conducted by maintaining the Si and C source gas ratioconstant during the entire crystal growth and by varying the dopantconcentration. The theoretical considerations and experimental datadiscussed in Kim do not give any insights into how the varying of the Siand C source gas ratio during crystal growth would affect dopantincorporation into the growing crystal. The study in Kim was limitedsolely to the effects on dopant concentration in grown crystals in thepresence of different partial pressures of dopant gas in the reactor.Kim also does not teach the varying of dopant gas in the reactor oncecrystal growth has begun. All the experiments and data collectedresulted from crystal growth based upon constant Si, C and dopant gasconcentrations. The varying of any gas within the reactor during crystalgrowth is not disclosed in Kim.

Another technique reported for InP and GaAs CVD growth systems are theuse of a blocking technique whereby relatively large amounts of acrystal compound are used to block or shield the crystal surface fromimpurities. Chemistry of The In--H₂ PCL₃ Process by R. C. Clarke, Inst.Phys. Conf.; Ser. No. 45; Chapter 1 (1979) 19-27 and Doping Behavior ofSilicon and Vapor Growth III-V Epitaxial Films by H. P. Pogue and B. M.Quimlich, J. Crystal Growth 31 (1975) 183-89. Although the blockingtechnique reduces contaminant incorporation into a crystal, the requireduse of large amounts of a crystal growing compound to affect blockinghas adverse effects on crystal quality and surface morphology. Inaddition to the problems associated with contaminant exclusion, thecontrolled intentional doping of crystals at relatively lowconcentrations has of yet been unachievable and/or unreproducible.Furthermore, techniques to control and/or reproduce sharp dopantconcentration profiles within crystals (i.e. from p-type to n-type,degenerate to lightly doped) are also unavailable.

In CVD crystal growth, control of the source gases in the prior art wassolely used for setting a single optimum source gas ratio for obtainingsmooth single-crystal surface morphology while preventing the formationof unwanted crystal formations (i.e. polycrystalline SiC in SiCcrystals), as well as other crystal defects. The teachings and theunderstandings in the art of CVD crystal growth taught against thechanging of the source gas ratio during crystal growth. It was believedthat one must first obtain the optimum source gas ratio in order toobtain high quality crystals. Once the source gas ratio set point wasobtained, dopant incorporation was controlled solely by controlling thedopant concentration in the reaction chamber. The prior art taught thatthe varying of the source gas ratio had no effect on the dopantincorporation (i.e. electrical properties) of the grown crystal. Oncedetermined, this optimum source gas ratio was constantly maintained andis specific and distinct for each CVD system during the growth of theentire crystal. The controlled constant amount of Si-source and C-sourcegases, taught by prior art, resulted in a constant ratio of availableSi-sites to C-sites (i.e., constant site composition) for incorporationof either crystal-growth atoms (i.e. Si and C atoms) or dopant-sourceatoms (e.g., Al, N, Ga, Ti, V, Na, Fe, P, S, O, B, F, etc.). The maximumconcentration of a p-type dopant incorporated into a growing SiC crystalis fixed, because the p-type dopant atoms compete with the Si-sourceatoms for a fixed amount of available Si-sites. When the amount ofaluminum dopant-source (Al) is increased in an attempt to obtain ahigher carrier concentration than this fixed amount (p=6×10¹⁷ cm⁻³), theexcess Al results in Al-containing material deposited on the reactorwalls which results in a very rough surface of a heavily p-doped(p=1.6×10¹⁹ cm⁻³) quasi-polycrystalline film. Therefore, it was notpossible to grow p-type doped (e.g. Al-doped) single crystal films withcarrier concentrations greater than p=6×10¹⁷ cm⁻³.

A variety of prior art techniques have been developed, and tried, in anattempt to control the impurity incorporation in the growing SiC films.For example, molecular beam epitaxy (MBE) utilizes an ultrahigh vacuumsystem and the growth of the crystal occurs by using a stream ofmolecules which is formed into a beam and focused onto a heatedsubstrate. This technique does offer some amount of control overdopant-concentration profiles. However, the low rate of growth is amajor drawback and poses several problems. First, commercialization isnot practical because of the very low growth rate. Second, the impurityincorporation is still limited by the purity of the source gas andcleanliness of the growth reactor. Furthermore, the problem of impurityincorporation is exacerbated by the very slow growth rate.

Another technique for doping SiC is the use of ion implantation, whichis a post-growth doping technique used to introduce the desired dopants.See Large-Band-Cap, III-V Nitride, and II-VI Zn--Se-Based SemiconductorDevice Technologies, by H. Morkoc et al; J. App. Phys. 76(3), p. 1368(1994). This method produces a large amount of damage to the crystalstructure and typically requires a post-anneal step to reduce the highdensity of defects (which greatly affects device quality) generated bythis technique. Furthermore, as a result of the unusually hightemperatures (>1800° C.) needed for the only partially effectiveannealing of SiC crystals, the dopant concentration changes by a factorof four (4×) for p-type and is lost via out-diffusion for the n-typedopants.

When degenerately doped layers are desired, such as the case when metalcontact layers are needed, obtaining degenerate p-type via CVD islimited by problems such as gas phase nucleation (from an excessivelyhigh concentration of p-type source gas needed) which results in verypoor film morphologies.

As stated above, silicon carbide (SiC) is emerging as a material ofchoice for fabrication of high power and/or microwave-frequencysemiconductor devices suitable for operation in high temperature, highradiation, and corrosive environments. The intrinsic material advantagesof silicon carbide are currently being exploited in the development ofhigh power and high frequency semiconductor devices for service in hightemperature, corrosive, and high radiation environments. Thesemicroelectronic devices include high voltage Schottky rectifiers andpower metal-oxide semiconductor field-effect transistors (MOSFETs),microwave and millimeter-wave devices, and high temperature, radiationresistant junction field-effect transistors (JFETs). However, in orderfor the theoretically calculated advantages of using SiC to be fullyrealized, advancements are needed both in the bulk growth and in theepilayer growth of SiC. For example, improvements in the bulk growth ofSiC are needed for elimination of device limiting micropores andmicropipes. Other advancements are needed in the epilayer growthprocess. In particular, dopant incorporation during the growth of SiCepilayers must be understood and reliably controlled. In order for theinherently superior high-temperature semiconductor properties of SiC tobe realized in advanced electronic device applications, control over theelectronic properties of chemical vapor deposited (CVD) spitaxial layersmust be improved. Prior to this work, control over dopant incorporationfor CVD SiC epilayers had been limited, with reproducible dopingtypically confined to doping concentrations ranging from N_(D) =2×10¹⁶cm⁻³ to 5×10¹⁸ cm⁻³ for n type and from N_(A) =2×10¹⁶ cm⁻³ to 1×10¹⁸cm⁻³ for p-type 6H--SiC epilayers. Expanding the reproducible dopingrange to include lower concentrations would enable the fabrication ofmultikilovolt SiC power devices, whereas the availability of higherdoping concentrations would result in devices with increasedperformances because of lower parasitic resistance.

As a result of the inadequate techniques available to control dopantand/or contaminant incorporation into crystals grown in a CVD process,there is a demand for a method to selectively exclude impurities from aCVD grown crystal.

THE INVENTION

The present invention relates to a method of growing high-qualitycrystals and controlling the impurity or dopant incorporation duringcrystal growth.

In accordance with the principal feature of the invention, there areprovided a system and a method of controlling the amount of a selectedelement deposited in a given growth area of a crystal grown by thechemical vapor deposition (CVD) process. One type of crystal grown is aSiC crystal wherein silicon (Si) is deposited in Si growth sites andcarbon is deposited in C growth sites at the growth area. The selectedelement competes for either the Si growth site or C growth site duringSiC crystal growth. The improvement comprises flowing of a first amountof a gaseous Si compound through the growth chamber and flowing a secondamount of a gaseous C compound through the growth chamber andcontrolling the ratio of the first amount of Si compound relative to thesecond amount of C compound to control the amount of the selectedelement deposited in the SiC crystal at the crystal growth area. Theselected element is an impurity and is either a dopant or a contaminant.

A crystal growth chamber is used to grow the crystals by chemical vapordeposition. The crystal growing chamber includes a crystal growing areaupon which crystals are grown. The crystal growing chamber may alsoinclude a heating element to control the temperature within the crystalgrowing chamber. The crystals are formed within the crystal growingchamber by introducing vaporized crystal growing compounds into thechamber. A carrier gas may be used to introduce the crystal growingcompounds into the crystal growing chamber. Many different types ofcrystals may be grown in the crystal growing chamber such as siliconcarbide, gallium arsenide, gallium phosphide, etc. The grown crystalsare formed by crystal components being deposited in crystal growingsites. For example, SiC crystals are grown by introducing a siliconcompound and a carbon compound into the crystal growing chamber. In thecrystal growing chamber, carbon atoms dissociate from the precursor ofthe carbon compound and silicon atoms dissociate from the precursor ofthe silicon compound. The SiC crystal is then grown by carbon atomsbeing deposited in carbon sites and silicon atoms being deposited insilicon sites. The SiC crystal is formed from stacked double layers orfilm layers of silicon and carbon atoms. Each film layer is formed froma layer of carbon atoms bonded to a layer of silicon atoms. Due to thedouble layer stacking of SiC film layers, the SiC crystal has a siliconface and a carbon face. The crystal growing sites, such as carbon sitesand silicon sites for a SiC crystal, are highly specific as to whattypes of atoms or molecules may enter the site during crystal growth.The high specificity of the crystal growing site is due in part to thephysical configuration of the site. Each atom and molecule has adistinct size and atomic configuration. Unless an atom or molecule havethe same or a similar size and configuration as the crystal growingsite, the atom or molecule will have a low probability of bonding at thegrowing site. The crystal growing site specificity is also affected bybonding forces in and about the growing site. Each atom and molecule hasan electron shell, having some charge, which interacts with electronshells of an atom or molecules present about the growing site, therebybeing either attracted to or repelled from the growth site. The purityof any crystal grown by the CVD process can be controlled by controllingthe impurity deposition in the growth sites during crystal growth.During crystal growth, a specific growth site has a particular demandfor a compatible atom or molecule. The growth site demand can bemanipulated by controlling the amount of compatible atoms or moleculesnear the growth site. The competition for a particular growth sitebetween a crystal atom component and an impurity atom or molecule iscontrolled by controlling the concentration of the respective atoms ormolecules at the growth site. This concept is herein referred to assite-competition epitaxy which serves to increase the reproducibility ofCVD crystal epilayers while greatly expanding the doping range for bothn-type and p-type CVD crystal epilayers. Surprisingly, the growthsite-competition can be accurately and reproducibly controlled bycontrolling the ratio of the crystal growing compounds introduced intothe crystal growing chamber. Many crystals other than SiC crystals whichhave two or more growth sites potentially can be grown by a CVD processand use the site competition technique to control impurityincorporation. The recognition and implementation of growthsite-competition technique allows for substantially greater latitude andcontrol as compared to prior growth techniques over crystal film growthconditions, so as to optimize crystal growth rate, surface morphology,impurity profiles and other film characteristics during crystal growth.

In accordance with a broader aspect of the present invention, there areprovided a system and a method of control ling the amount of anon-crystal element deposited in a given growth area of a crystal formedfrom at least two crystal elements as the crystal is grown at thecrystal growing area by a CVD process conducted in a growth chamber. Thecrystal comprises a first crystal element and a second crystal elementand the grown crystal has at least two crystal growing sites. The firstcrystal element is deposited in the first crystal growing site and thesecond crystal element is deposited in the second crystal growing site.The non-crystal element is competitive for at least one of the growthsites. The improved method comprises flowing a controlled amount ofgaseous crystal element compounds through the growth chamber, whereineach of the crystal element compounds includes at least one of thecrystal elements, and controlling the ratio of the controlled amount ofgaseous crystal element compounds to control the amount of thenon-crystal element deposited in a particular growth site of the crystalat the crystal growth area. The non-crystal element is an impurity suchas a dopant or a contaminant. The crystal can be formed by two, three,four or more crystal elements to form a binary, tertiary, etc., type ofcrystal. Each crystal element is deposited in a particular crystalgrowth site. The non-crystal element competes with at least oneparticular crystal growth site. The non-crystal element normally onlycompetes with one crystal growth site due to the geometry and charge ofthe non-crystal element and the geometry and charges of the crystalgrowing sites. If a non-crystal element can deposit itself in more thanone crystal growing site, the non-crystal element usually has a muchlarger affinity for one over the others. The types of crystals thatpossibly can be grown include ZnSe, GaAs, InP, In_(X) Ga_(1-X) As,GaAsP, GaP, InAs, In_(X) Ga_(1-X) As_(Y) P_(1-Y), etc.

In accordance with another aspect of the present invention, a dopantmaterial is introduced into the crystal growing chamber and the rate ofdopant incorporation into the crystal during crystal growth iscontrolled. Various types of dopant materials may be used to form n-typeor p-type layers in the crystal. The dopant profile of a crystal is veryimportant for the type of electronic device the doped crystal isincorporated therein. The dopant material is introduced into the growingchamber either in its pure form (i.e. nitrogen) or as a compound whichincludes the dopant material bound to a dopant precursor (i.e. trimethylaluminum). The dopant material which is introduced into the growthchamber is selected such that the dopant material has an affinity for atleast one of the growth sites of the crystal. Almost all dopants have anaffinity for only one growth site due to the physical characteristics ofthe dopant and of the growth site. The select few dopants which may bindin multiple crystal growth sites have a higher affinity for one growthsite as compared to the others. Once the dopant material is introducedinto the crystal growing chamber, the dopant material primarily competeswith a crystal element component for a particular growth site. Thecompetition (i.e. demand) of the dopant material with a crystal atomcomponent for a particular growth site is controlled by the manipulationof the availability of the growth site. The manipulation of the growthsite demand is accomplished by adjusting the ratios of the crystal atomcomponents within the crystal growing chamber. By proper manipulation ofthe ratios of the various crystal atom components, a particular growthsite will become more or less available to the dopant material duringthe crystal growing period, thus affecting and controlling the growthsite demand and the dopant concentration profile within a particularcrystal.

In accordance with still another aspect of the present invention, thecrystal is grown on a pretreated substrate to expand the growingparameters for producing high quality crystals. The substrate may or maynot be made of the same material as the grown crystal film. A crystalfilm grown on a substrate may be grown homoepitaxially (substrate andcrystal film have the same crystal elements and structure) or grownheteroepitaxially (substrate and crystal film have different crystalelements and/or structure). The substrate is pretreated to removeimpurities from the substrate surface which may cause defective crystalfilms during the crystal growth. Abnormal crystal growth can take placeat sites where there is contamination or defects or some other surfacedisturbance on the substrate surface. These contaminants or defectsresult in unwanted nucleation resulting in crystals having unwantedpolytype structure, inferior surface morphology, stacking faults, APBs,low quality dopant profiles, high contaminant concentrations, etc.Contaminants and surface defects are removed from the substrate byproper cutting and polishing of the substrate surface and subsequentetching of the substrate surface. The etching is carried out so as notto alter the substrate surface in a manner that would impair crystalgrowth upon the substrate surface. The clean, low-defect substratesurface surprisingly expands the operation parameters for growingquality crystals. By reducing the stress to the crystal growing process,previous Si/C concentration ratios within a growth chamber were limitedin range and did not stray far from a narrow ratio-range which isdependent on the choice of a hydrocarbon employed (Si:C:1:1: for SiH₄/C₂ H₄, Si:C::2:3 for SiH₄ /C₃ H₈). With any given crystal growingprocess, there are stress components to the process such as temperature,pressure, crystal component concentrations, dopant concentrations,impurity concentrations, growth surfaces, etc., which affect the crystalfilm layer during growth. The elimination or reduction of stressesduring crystal growth caused by a defective and/or contaminatedsubstrate surface allows for a greater variance in stress levels fromother stress components. The reduced stresses caused by a pretreatedsubstrate enable the use of greater crystal compound ratio ranges in thegrowth chamber for enhanced control of impurity incorporation into thecrystal.

In accordance with still another aspect of the present invention, thereis provided a method for growing crystals whereby the crystals are grownon a particular surface face of a substrate to reduce the stresses tothe crystal growing system. A substrate formed from two or moredifferent atoms will have a crystal structure having at least threedistinctive faces. Substrates formed of Si--C crystals have a siliconface (Si-face), a carbon face (C-face) and a carbon-silicon face(A-face). Each face of the crystal structure has different physicalcharacteristics (e.g. polarity) which can increase or reduce thestresses to the crystal growing system, which will affect the quality ofthe crystals grown on the substrate face.

The primary object of the present invention is the provision of a methodto grow high-quality, reproducible crystal films by chemical vapordeposition.

Another object of the present invention is the provision of a method forcontrolling impurity incorporation in a crystal film layer during thegrowth of a crystal.

Still another object of the present invention is the provision of amethod to grow high-quality, low-defect reproducible crystals having adesired impurity profile.

Yet still another object of the present invention is to provide a methodfor growing crystals on a pretreated substrate whereby contaminants andsurface defects have been removed from the surface of the substrate uponwhich crystals are to be grown.

Another object of the present invention is to provide a method forgrowing crystals on a particular face of the substrate to produce highquality crystals having low defects and a smooth surface morphology.

Yet another object of the present invention is to provide a method fordegenerate doping of crystals for both p-type and n-type epilayershaving a high-quality surface morphology and low defects.

Yet still another object of the present invention is to provide a methodfor growing crystals having very abrupt changes in dopant n-p typeprofiles.

Another object of the present invention is to provide a method ofgrowing extremely abrupt dopant profile interfaces.

Yet another object of the present invention is to provide a method forgrowing low-defect, low-concentration doped single-crystals on varioussubstrates.

Another object of the present invention is to provide a method forgrowing crystals having abrupt changes in dopant concentration profiles.

Still yet another object of the present invention is to provide a methodfor growing high-quality, low-defect crystals having gradient dopantprofiles.

Yet another object of the present invention is to provide a method forincreasing the range of concentration ratios usable to growhigh-quality, low-defect crystals.

Yet still another object of the present invention is the provision of amethod to grow high-quality, low-defect crystals having very high dopantconcentrations.

These and other objects and advantages will become apparent to thoseskilled in the art upon reading the following description taken togetherwith the accompanying drawings.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic drawing of a CVD system employed for the growingand doping of crystals in accordance with the present invention;

FIG. 2 is a plan view of a substrate that has been divided into selectedgrowth regions;

FIG. 3 is a plan view schematic drawing of a selected growth regionhaving parallel lines which illustrate some of the atomic-scale growthsteps;

FIG. 4 is a schematic diagram illustrating the homoepitaxial growth of6H--SiC films upon a 6H--SiC substrate;

FIG. 5 illustrates a plan view of a 6H--SiC crystal showing thedimensional axis of the crystal;

FIG. 6 is a cross-sectional view of a SiC substrate showing the Si-faceand the C-face and the A-face of the substrate;

FIG. 7 is a cross-sectional view of a 6H--SiC substrate nucleated with3C--SiC;

FIG. 8 is a substrate as shown in FIG. 7 illustrating the expansion of3C--SiC growth from a nucleation point on the 6H--SiC substrate;

FIG. 9 illustrates the site competition at the C site and Si sitesduring the growing of a SiC crystal;

FIG. 10 is a graphic illustration of a dopant concentration profile of acrystal;

FIG. 11 is a cross-sectional view of p and n epilayers of a SiC crystalgrown upon a SiC substrate; and

FIG. 12 illustrates the forward and reverse current-voltagecharacteristics on a logarithmic scale of a 3C--SiC crystal at severaltemperatures.

PREFERRED EMBODIMENT

Referring now to the drawings, wherein the showings are for the purposeof illustrating preferred embodiments of the invention only and not forthe purpose of limiting the same, the invention describes an improvedchemical vapor deposition (CVD) method for obtaining improved control ofimpurity incorporation and dopant profiles of CVD films and alsoimproving the quality of the grown crystal films. While the method maybe applied to many different crystals, the method will be specificallydescribed with respect to the growing of silicon-carbide (SiC) crystals.The improved CVD method includes pretreating a substrate, heating thesubstrate in a reaction chamber, introducing a carrier gas, vaporizingthe crystal growing compounds, introducing the vaporized compounds inthe reaction chamber via the carrier gas, maintaining proper energylevels and material flow rates in the reaction chamber for a sufficienttime to grow a crystal film having a desired smooth surface morphology,a uniform thickness, a low-defect density and a controlled impurityprofile. The crystals may be intentionally doped to form n-type and/orp-type crystals. The improved CVD crystal growing method is based on thediscovery that impurity incorporation into a growing SiC film layers isvery sensitive to the ratio of silicon compound (Si containingprecursors) and carbon compound (C containing precursors) in thereaction chamber during crystal growth. By varying the Si/C compoundratio at or near the crystal growing surface, the impurity incorporationinto the growing crystal film is controlled.

The method of the invention can be carried out with a conventionalchemical vapor deposition (CVD) system similar to that used in Si, SiCand GaAs semiconductor technology. The gases used in a SiC CVD systemare hydrogen (used as a carrier gas), Silane (used as a source of Si),HCL (used for cleaning and etching the substrate surface), propane (usedas a source of C), nitrogen (N₂) (used as a n-type dopant), andtrimethyl aluminum (TMA) (used as a p-type dopant). Other gases may beused as the Si or C source or used to dope the crystal. If organiccompounds are used as the Si and C source, the process is commonlyreferred to as metal-organic vapor phase epitaxy (MOVPE). Any CVD systemthat can deliver these gases to a suitable reaction chamber at theproper flow rates under high purity conditions can be used for theinventive method.

Referring now to FIG. 1, there is shown a schematic, partial view of asuitable CVD reaction system for carrying out the process of theinvention. The CVD reaction system includes a reaction chamber 22comprising a double-walled quartz tube such that the inner quartz tubecan be water cooled. The inside diameter of reaction chamber 22 ispreferably 50 mm. A SiC substrate 24 is supported by a SiC coatedgraphite susceptor 26, which in turn is supported by quartz support 28.To produce the desired temperature of the surface of substrate 24, aradio-frequency (RF) induction coil 30 is disposed around reactionchamber 22. Induction coil 30 is powered by frequency generator 31. TheRF field produced by induction coil 30 heats substrate 24 via susceptor26 to the desired temperature. When the SiC film layers are grown,substrate 24 is preferably a SiC substrate. The gaseous crystalcompounds are introduced into reaction chamber 22 by primary line 33.Primary line 33 is located at one end of reaction chamber 22 and directsthe gases to flow in direction G across substrate 24 and out theopposite end of chamber 22. The various gaseous crystal compounds areconnected to primary line 33 and the gas flow is regulated by valves 34and regulators 35 connected to each gas line. Line 36 is the silicon gasline that controls the silane flow into primary line 33, and line 37 isthe carbon gas line that controls the propane flow into primary line 33.The dopants are introduced into primary line 33 by line 38 and line 39.Line 38 is the n-type dopant line and preferably controls the nitrogengas (N₂) flow rate. Line 39 is the p-type dopant line and preferablycontrols the trimethyl aluminum (TMA) flow rate. Carrier gas line 31carries all the gaseous crystal compounds and dopants through primaryline 33 and into reaction chamber 22. The carrier gas is preferably agas such as hydrogen gas (H₂). Carrier gas line 31 is partially divertedinto line 31a to supply line 39 so that the carrier gas can be bubbledthrough the liquid TMA. A vacuum line V connected to a vacuum can beconnected to primary line 33 to evacuate reaction chamber 22 of gases.

Preferably, substrate 24 is pretreated to remove any contaminants orimpurities on the surface of the substrate so as to facilitate thegrowing of high-quality, low-defect epitaxial films. SiC substrate 24 isprepared by slicing a section from a SiC boule. Substrate 24 may be cutsuch that the surface is slightly misoriented relative to the basalplane by some tilt angle. If 3C--SiC film layers are to beheteroepitaxially grown on an α-SiC substrate 24, the tilt anglepreferably is less than 1° and additional surface preparation isnecessary. If SiC film layers are to be grown homoepitaxially on SiCsubstrate 24, the tilt angle is preferably greater than 0.1°. The tiltdirection is preferably toward the <1100> or <1120> direction, asillustrated in FIG. 5, to produce the optimum growth rates and qualityof the SiC epitaxial films grown on substrate 24. The surface ofsubstrate 24 is polished preferably on one surface with a diamond paste.SiC substrate 24 has three faces, a Si-face 50, a C-face 52 and theA-face 54, as illustrated in FIG. 6. Any of the faces may be polishedand prepared for growth of the SiC epitaxial layers. Preferably, Si-face50 is polished and used for epitaxial growth. It has been found thatSi-face 50 produces the highest-quality epitaxial layer films which havethe best surface morphology and lowest defects.

If heteroepitaxy is preferred, then substrate 24 is further preparedlycreating boundaries or grooves 62 on the face of substrate 24 which formgrowth regions 60, as illustrated in FIG. 2. Grooves 62 forming growthregion boundaries 60 are preferably cut with a precision dicing saw witha 25 micrometer thin blade to minimize crystal damage; however,boundaries 60 may be formed by other processes such as photolithography,ion etching and/or photochemical or electrochemical etching processes.The width of groove 60 need only be less than 1 micrometer but largerwidths can also be used. The depth of groove 60 is preferably about 50micrometers but may be larger or smaller.

Once the substrate surface has been polished and growth regions 62 havebeen formed, substrate 24 is placed in reaction chamber 22. Prior togrowing the crystal film layers on substrate 24, the substrate ispretreated with a pregrowth etch to remove contaminants and defects onthe surface of the substrate that could act as unwanted sites forheterogeneous nucleation of the SiC film layers. These defects on thesurface of the substrate can be generated during the cutting andpolishing of the substrate. Preferably, the pregrowth etch involvessubjecting substrate 24 to a high temperature gaseous etch in a mixtureof hydrogen chloride gas and hydrogen within the reaction chamber. Theetch is monitored such that the substrate is not altered in a way thatunwanted sites for heterogeneous nucleation are introduced to thesurface of the substrate. Preferably, the etch uniformly removes atleast one atomic layer from the surface of substrate 24 to ensure alow-defect, highly-pure surface. A typical etch is carried out for about25 minutes at a temperature of 1350° C. using about 3-4% hydrogenchloride gas in an H₂ carrier gas with a flow of about 3 liters perminute. Preferably, the concentration of the hydrogen chloride gasranges between 1-5% during the pregrowth etch. Lower hydrogen chloridegas concentrations may not properly remove all the contaminants andsurface defects from the substrate. Higher hydrogen chloride gasconcentrations may produce a rough surface morphology or pits on thesubstrate, which may cause lateral growth of the epitaxial layers andcreate random nucleation sites throughout the surface of the substrate.The temperature during the etch ranges between 1200°-1500° C. Lowertemperatures will not properly eliminate unwanted heterogeneousnucleation sites. Temperatures greater than 1500° C. will too rapidlyetch the substrate surface around the peripheral edge of the substrateand introduce unwanted heterogeneous nucleation sites upon the surfaceof the substrate. Other pregrowth treatments, such as oxidation orreactive ion etching, may also be used to further remove potentialunwanted nucleation sites prior to growing the crystal epilayers.

Referring now to FIG. 3, nucleation sites 70 on substrate 24 may beformed for intentional heterogeneous nucleation. Nucleation sites 70 canbe formed by an intentional localized alteration of the surface ofsubstrate 24. These sites 70 can be formed by indenting substrate 24with a diamond scribe at a predetermined location, damaging substrate 24by using an electrical probe striking an arc between the surface, and/orimplanting a single-crystal whisker of some desired polytype onsubstrate 24. If 3C--SiC crystals are to be grown on a 6H--SiC substrate24, as shown in FIG. 8, a whisker of 3C--SiC polytype crystal isimplanted at nucleation site 70 in a desired growth region 62 ofsubstrate 24, as illustrated in FIGS. 3 and 7. For growing SiCepilayers, the optimum location for nucleation site 70 is at the cornerof growth region 62 and on the topmost terrace or step 72, asillustrated in FIG. 3. As illustrated in FIGS. 3 and 7, the surface ofsubstrate 24 comprises several steps 72 of crystal layers. The multiplesteps 72 are formed when the substrate is cut at any non-zero anglerelative to the basal plane. For instance, if the substrate surface istilted at an angle α of 3° (relative to the basal plane), as illustratedin FIG. 4, a substrate surface will cut across several crystal layers toform multiple crystal layer steps 72. Nucleation site 70 should belocated on the uppermost step 72, as illustrated in FIG. 3. The optimumdirection of the substrate tilt is along the diagonal D of growth region60, as illustrated in FIG. 3.

Once substrate 24 has been pretreated, reaction chamber 22 is preparedfor crystal growth. Reaction chamber 22 is preferably evacuated byvacuum via vacuum line V and subsequently purged with an inert gas toremove impurities. Hydrogen gas may be used to purge the reactionchamber. Once the reaction chamber is purged, the carrier gas flow ratesand the temperature within the reaction chamber are brought toequilibrium. Hydrogen gas is preferably used as the carrier gas, butother gases (e.g. inert gases) can be used. Once the temperature andflow within the reaction chamber have reached equilibrium, generallywithin less than one minute, silane and propane are added to the carriergas to initiate SiC growth. Preferably, the silane concentration withinthe carrier gas is approximately 200 ppm resulting in a 200 ppm atomicconcentration of Si. The amount of propane introduced into the carriergas is approximately 130 ppm to 600 ppm resulting in a atomicconcentration of C between 390 ppm to 1800 ppm. The prescribedpretreatment of substrate 24 allows for significantly greater deviationsfrom the optimum Si/C ratio than was previously thought possible forgrowing high-quality, low-defect SiC crystals. The ratio of the atomicconcentrations of Si to C may be varied to create different growth ratesand different types (i.e. n or p-type) of SiC epilayers. The ratio mayrange between 0.01-1.0 and preferably is between 0.1-0.5.

Referring now to FIG. 4, there is shown an atomic-scale cross-sectionaldrawing of 6H--SiC substrate 24 comprising several layers of 6H--SiCfilm 41 and several epitaxial 6H--SiC film layers 40 deposited on thesurface of the substrate. SiC epilayer growth rates from a carrier gascontaining 200 ppm silane and 600 ppm propane resulted in a verticalepilayer film growth rate parallel to the c-axis of about 5.5micrometers per hour. Once the crystal begins to grow, multiple layers40 will form on top of each other thus producing a multiple-layer SiCepitaxial film, as shown in FIG. 4.

FIGS. 7 and 8 illustrate the growth of 3C--SiC epilayers 40a onsubstrate 24 made of 6H--SiC epilayers 41. As the crystal growthcontinues over time, the 3C--SiC epilayers 40a grow laterally from3C--SiC nucleation site 70a until the 3C--SiC epilayers (film layers)40a completely cover the growth region. The growing of 3C--SiC boundaryfilm layer 40a on 6H--SiC film layer 40 is called a heteroepitaxiallayer. A heteroepitaxial layer is an epitaxial film layer that is of adifferent material, different polytype, or is lattice-mismatched fromthe film layer upon which it grows. The heteroepitaxial layer is understress due to compression or tension along the plane between the twodifferent polytype epilayers. Although there is stress between the twopolytype epilayers, the 3C--SiC film layers 40a have few, if any, DPBsand stacking faults because nucleation of the 3C--SiC epilayer takesplace at one location.

Referring now to FIG. 9, during the growth of the SiC epilayers Si atomsare deposited in Si sites 80 and C atoms are deposited in C sites 82. Asa result, the SiC epilayers are formed by layers of Si atoms 86 and Catoms 84 forming on substrate 24. One epilayer or film layer isrepresented as one double-stack layer of a Si atom layer 86 and C atomlayer 84. The SiC epilayers grow in stacking alternating layers of Silayers 86 and C layers 84, which increase the thickness of the SiCcrystal. The type of SiC epilayer grown on the substrate can becontrolled by nucleation sites and/or the polytype of the substrate. Asillustrated in FIG. 4, 6H--SiC epilayers 40 are grown on a 6H--SiCpolytype substrate 24, which is referred to as homoepitaxial growth. Asillustrated in FIG. 8, 3C--SiC epilayers 40a are grown on a 6H--SiCpolytype substrate using 3C--SiC nucleation site 70a, which is known asheteroepitaxial growth.

FIG. 9 is an atomic scale schematic illustration of Si atom layers 86and C atom layers 84 forming on substrate 24. Si-face 50 is prepared forSiC epilayer growth resulting in C atom layer 84 being the first atomlayer to form on the surface of the substrate. The growth of C atomlayer 84 on Si-face 50 results in the growth of the highest-quality,lowest-defect SiC epilayers. During the growth of the SiC film layers,there are a number of different types of atoms or molecules at or nearthe film surface. H₂ represents the carrier gas in the reaction chamber.C_(Y) and Si_(X) represent a C atom and a Si atom respectively whichhave dissociated from their respective precursors and are in the processof depositing themselves in their respective C sites or Si sites. X, Yand Z represent impurities which intentionally or unintentionally existin the reaction chamber. Specifically, X represents a dopant thatcompetes with an Si_(X) (Si atom) for Si site 82 and Y represents adopant that competes with a C_(Y) (C atom) for C site 80. Z is acontaminant that was not removed during the purging of the reactionchamber or was unintentionally introduced into the reaction chamber. Zmay compete for C site 80, Si site 82, both sites or no site.

During the growth of the SiC film layer, the amount and type of impurityincorporated into the film layer is controlled to produce a SiC crystalwith specifically designed properties. Dopants can be intentionallyadded to alter the electrical and/or optical characteristics of the SiCcrystals. Phosphorous and nitrogen dopants can be added to the SiC filmlayers to form n-type layers and aluminum and boron can be added to SiCfilm layers to form p-type layers. When an n-type SiC epilayer isformed, phosphorous or nitrogen, as represented as Y in FIG. 9, competeswith C_(Y) for C sites 80 during the formation of the C atom layer 84.Similarly, p-type dopants such as aluminum and boron, as represented byX in FIG. 9, compete with Si_(X) for Si sites 82 during the formation ofthe Si atom layer 86. The impurity incorporation into each atom layer issignificantly affected by the ratio of Si/C in the reaction chamber.During the growing of a p-type film layer, wherein aluminum, sodium,iron, gallium, titanium, vanadium, or boron is incorporated into Si atomlayer 86, the amount of p-type dopant incorporated is increased bydecreasing the Si/C ratio in the reaction chamber. Conversely, theamount of p-type dopant incorporated into Si atom layer 86 is decreasedby increasing the Si/C ratio in the reaction chamber. Similarly, when an-type film layer is to be grown, wherein oxygen or nitrogen isincorporated into C atom layer 84, the amount of n-type dopantincorporation increases as the Si/C ratio increases and n-type dopantincorporation decreases as the Si/C ratio decreases. The inventors havedetermined that for n-type doping, fluorine, chlorine, sulfur andnitrogen incorporate into Si-sites; and oxygen and nitrogen incorporatein C-sites. For p-type doping, sodium, iron, aluminum, gallium,titanium, vanadium, and boron incorporate into Si-sites, and boronincorporates into C-sites.

Site-competition epitaxy is an advancement for the control of dopantincorporation such as for both p-type and n-type doped epilayers, whichhas resulted in increased doping range and improved dopingreproducibility for the growth of chemical vapor deposition (CVD) SiCepilayers. Use of site-competition epitaxy will lead to improved deviceperformance which includes high voltage diodes, ohmic as-depositedcontacts, and high temperatures JFETs. The usefulness ofsite-competition epitaxy relies upon each particular dopant atomsubstituting primarily for either a Si atom in the Si-site or for a Catom in the C-site of the growing SiC epilayer. Site-competition epitaxyhas been successfully used for control of both Al and N doping, partlybecause Al substitutes for Si in the Si-sites whereas N mainlysubstitutes for C in the C-sites of the SiC lattice when grown of theSi-face. The silicon-to-carbon (Si/C) ratio within the growth reactorstrongly affects dopant incorporation for epilayers grown on the(0001)6H--SiC silicon face (Si-face), SiC epitaxy on SiC(1210) A-facesubstrates and on 3C--SiC(111) and 4H--SiC(0001) Si-face basal planesubstrates and on SiC (0001) C-face substrates.

The site-competition epitaxy process is based on the discovery that theconcentration of dopant atoms incorporated into a growing SiC crystal isvery sensitive to the Si-source/C-source gas ratio. That is, theconcentration of dopant incorporated into a growing SiC crystal can becontrollably varied solely by purposely changing the Si/C ratio for eachselected single amount of a dopant source in the reaction chamber. Theuse of site-competition epitaxy affords production of very high qualityp-type SiC single crystal films including carrier concentrations withp>5×10¹⁸ cm⁻³. In fact, concentrations at least as high as p>5×10²⁰ cm⁻³have been obtained. The site-competition epitaxy process provides forgreater incorporation of p-type such as Al by decreasing theSi-source/C-source ratio. This results in 1) a decreased amount ofSi-source to compete with the p-type dopant source for availableSi-sites, and 2) an increase in the amount of available Si-site (i.e.,from increasing the ratio of available Si-site to C-sites) for enhancedp-type dopant incorporation. This is not possible by using any prior artteachings wherein a fixed Si/C ratio was used because theSi-source/C-source is maintained constant and therefore the relativeamount of Si-site is also to be constant. Analogously, films of muchlower doping concentrations can be produced by using site-competitionepitaxy. For example, Al doping levels can be lowered by increasing theSi-source/C-source ratio, thereby decreasing the relative amount ofavailable Si-sites while simultaneously increasing the amount ofSi-source to out compete the Al-dopant source for available Si-sites,both of which favor exclusion of Al from the growing SiC film. Thesuperior doping ranges for n-type doping are also obtained by usingsite-competition epitaxy to grow single crystal films including carrierconcentrations with n>2×10¹⁹ cm⁻³. N-type concentrations as high as atleast about n>5×10²⁰ cm⁻³ have been obtained. In addition, the Si/Cratio can be changed and/or selected at a single ratio to form lowdopant concentrations for both p-type and n-type dopants in singlecrystal films which are less than 1×10¹⁶ cm⁻³ and have been reproduciblyobtained at dopant levels as low as less than about 8×10¹³ cm⁻³. Thesehigh dopant concentrations (i.e.>5×10¹⁸ cm⁻³) and low dopantconcentrations (i.e.<2×10¹⁶ cm⁻³) were previously unachievable in grownsingle crystal films. High dopant concentrations could be achieved onlyby further crystal processing (i.e. ion implantation). However, suchpost growth crystal processing destroyed or damaged the crystal andfurther did not create a uniform dopant distribution in the crystal.

Although the inventors do not want to be held to one theory for thisphysical phenomena, it is believed that the controlled incorporation ofdopants into the SiC film layers is accomplished by the manipulation ofsite competition at C sites 80 and Si sites 82. Although dopantconcentrations in the film layers can be controlled somewhat byregulating the amount of dopant introduced into the reaction chamber,very low dopant concentrations, very high degenerate dopantconcentrations, sharp p-n junctions or n-p junctions, and/orreproducible dopant concentrations are unattainable by solely regulatingthe dopant concentration. The additional degree of control over dopantincorporation is carried out by manipulating the demand of a particulardopant at Si site 82 or C site 80 during the film layer growth. Duringthe growth of each SiC film layer, the rate at which a layer of Si atoms86 and C atoms 84 are formed depends on the availability of atoms and/ormolecules that can fill a particle site in the Si atom 86 or C atom 84layer. When the concentration of C atoms in the reaction chamber isincreased relative to the Si atom concentration, the demand for atoms tofill C sites 80 decreases since the available amount of C atoms hasincreased. As a result, the C atom layer 84 is formed at a faster rateand C sites 80 are disproportionately filled with C atoms. However, therelative increase in C atoms effects a relative decrease in Si atomsavailable to fill Si sites 82. The relative decrease in available Siatoms results in a slower filling of Si sites 82 and a greater demandfor any type of atom or molecule to fill Si site 82. The greater Si sitedemand results in an increase in non-Si atoms (i.e. dopants,contaminants) filling Si sites 82 to form Si atom layer 86. Conversely,a relative increase in Si atoms to C atoms in the reaction chamberdecreases the demand of atoms or molecules at Si sites 82 and increasesthe demand of atoms or molecules at C sites 80. The decrease in Si site82 demand, due to a relative increase in Si atom availability, resultsin a disproportionately high Si atom occupation in Si atom layer 86.Further, the increase in C site 80 demand reduces the rate at which theC atom layer 84 is formed and increases the amount of impuritiesrelative to C atoms incorporated into C atom layer 84.

The simultaneous control of the multiple growth sites of the crystal canbe used to control the impurity profile of crystals irrespective of theimpurities contained within the reaction chamber. For instance, duringthe growth of SiC crystals, a n-type crystal can be grown even if equalamounts of n-type and p-type dopant were present in the reactionchamber. By properly controlling the Si/C ratio (i.e. increasing theSi/C ratio) in the reaction chamber, the simultaneous demandmanipulation of Si site 82 and C site 80 is affected whereby Si site 82demand decreases and C site 80 demand increases. The increase in C site80 demand results in larger n-type dopant incorporation into C atomlayer 84 than p-type dopant incorporation into Si atom layer 86, therebyforming a n-type SiC crystal. The simultaneous growth site manipulationmay be possibly used to grow other types of crystals having at least twogrowth sites.

The phenomena of controlling impurity incorporation are also believed tobe affected by growth site physical properties, site geometry and thegrowth rate of Si atom layer 86 and C atom layer 84. It is postulatedthat the competition for C site 80 between a C atom and an impurity isin favor of the C atom, and the competition for Si site 82 between a Siatom and an impurity is in favor of an Si atom. In a crystal latticestructure, a particular growth site has a specific geometry into whichan atom or molecule can bond. For instance, in Si site 82, the geometryis of the shape of a Si atom. The electron cloud about the Si atomspecifically bonds with the atoms surrounding Si site 82 to form ahighly stable bond. An impurity which does not have the proper geometryto fit into Si site 82 has a very low probability of bonding. Si site 82also has a particular charge, due to atoms located about Si site 82,which is receptive to a Si atom or another atom or molecule having acharge similar to a Si atom. Furthermore, there exists an equilibriumreaction between the bonding and unbonding of a particular atom to Sisite 82. It is believed that during relatively rapid Si atom layergrowth, wherein the Si/C ratio is large, the equilibrium of reaction isin favor of the Si atom. In addition, it is believed that theequilibrium of reaction in favor of the Si atom reduces as the growthrate of the Si atom layer decreases. Therefore, in a situation wherein aSi atom and an impurity simultaneously reach Si site 82, the Si atom hasa higher probability to bond at Si site 82 due to its geometry, chargeand equilibrium of reaction with Si site 82. The postulated competitionfor Si atom in Si sites 82 equally applies to C atom competition in Catom sites 80.

The control of the Si/C ratio in the reaction chamber enables theproduction of quality, reproducible SiC crystals which can bedegenerately doped with dopant concentrations exceeding 1.0×10¹⁹ cm⁻³ orvery light dopant concentrations at least as low as 8.0×10¹³ cm⁻³.

FIG. 10 illustrates a dopant profile of an n-p type SiC crystal. The SiCcrystal was initially doped with a n-type dopant to produce severaldegenerate n-type film layers 90 containing n-type dopant concentrationsexceeding 1.0×10¹⁹ cm⁻³. The degenerate layers generally are used asconnect points between the crystal and metal connects (Ohmic contacts).The dopant profile illustrates SiC crystal film layers 92 containingdecreasing amounts of n-type dopant resulting from a decrease in dopantatoms and/or a decrease in the Si/C concentration in the reactionchamber. A sharp n-p 94 junction is illustrated. Junction 94 can beproduced by decreasing or terminating the n-type dopant flow into thereaction chamber and introducing or increasing p-type dopant into thereaction chamber. Further, the regulating of the Si/C concentrationratio in the reaction chamber can produce a highly defined, controllableand reproducible n-p junction. The p-type dopant incorporation isincreased in film layers 96 by increasing the dopant concentrationand/or decreasing the Si/C concentration ratio. The dopant incorporationis increased until p-type degenerate film layers 98 are formed.Alternatively, a p-n junction can be formed by first forming p-type filmlayers and then forming n-type film layers.

FIG. 11 illustrates a SiC crystal having p-type SiC layer 100 growing onthe surface of substrate 24, a n-type layer 102 grown on layer 100 and ap-type layer 104 grown on layer 102 to form a pnp SiC crystal which isthe basic building block of all electronic devices. Many different SiCcrystals with varying dopant profiles can be made by the inventivemethod. Controlled and reproducible dopant concentration below 1.0×10¹⁶cm⁻³ and degenerate film layers exceeding 1.0×10¹⁹ cm⁻³ (resulting in"Ohmic as deposited" metal contacts), which were previouslyunattainable, can be formed by growth site manipulation through thecontrol of Si/C concentration in the reaction chamber. The SiC filmlayers can be grown on a substrate surface with little or no tilt angleor with large tilt angles. The control of growth site competition byvarying the Si/C concentration in the reaction chamber appears not to beaffected (within reasonable limits) by the tilt angle of the surface ofthe substrate. However, the tilt angle has been found to affect thepolytype and crystal quality of the grown SiC crystals.

The purity of a crystal grown by a CVD process can be substantiallycontrolled by the manipulation of site competition at particular growthsites of the SiC film layers. During crystal growth, contaminants,illustrated as Z in FIG. 9, can incorporate themselves into the SiCcrystal during crystal growth. Purging the reaction chamber with anultra-purified gas, such as hydrogen, reduces some but not all thecontaminants within the reaction chamber. The source of the contaminantscan be the precursors for the crystal atoms, impurities whichinadvertently leak and/or are introduced into the reaction chamberand/or atoms or molecules within the reaction chamber prior to crystalgrowth. Preferably, the crystal component precursors are selected so asnot to behave as a contaminant source. A common contaminant in SiCcrystal growth is aluminum. Aluminum is commonly used as a p-typedopant. Upon entering the reaction chamber, aluminum, due to itsmetallic properties, has a tendency to remain in the reaction chamberafter the doping process has ended. During the growth of a pure SiCcrystal, aluminum remaining within the reaction chamber can bond in Sisites 82 and potentially degrade the electrical characteristics of theSiC crystal. The amount of aluminum contaminant incorporation into theSiC crystal can be significantly reduced by increasing the Si/Cconcentration ratio in the reaction chamber and reducing the Si growthsite demand. By identifying the primary contaminant which isincorporated within a SiC crystal and the growth site at which thecontaminant bonds, the reduction of contaminant incorporation in the SiCcrystal can be controlled by properly adjusting the crystal growingcompound concentration ratio within the reaction chamber to reduce thedemand of the contaminant at the growth site.

The inventive method can be applied to the fabrication of semiconductordevice structures of many kinds. The technique for control over thedoping is far superior to any known conventional doping technique forgrowing SiC crystals. The technique also provides for abrupt in situdoping during crystal growth and is capable of fully utilizing thedoping concentration ranging from degenerately p-type doped (for "Ohmicas deposited" contacts) through near intrinsic (extremely low dopedn-type or p-type) up to and including degenerate n-type SiC. The abilityto control the dopant in such a predictable and reproducible manner iscrucial in the fabrication of most SiC devices. Many modifications ofthe inventive method are possible; for example, the inventive methodcould be carried out in an ultrahigh vacuum system (e.g. a molecularbeam epitaxial (MBE) system).

EXAMPLE 1

A n-type SiC crystal containing about 1.0×10¹⁵ cm⁻³ was formed using theimproved CVD method. For the method described in this example, acommercial 6H--SiC substrate cut from a 6H--SiC boule was used. TheSi-face (0001) of the substrate was polished with a diamond paste andcut with a 25 mm dicing saw to form boundaries on the substrate surface.The substrate was subsequently placed within the reaction chamber andsubjected to an etch for about 20 minutes at a temperature of 1375° C.using about 3%-4% HCL gas in a H₂ carrier gas with a flow of about 3L/min. After 25 minutes, the flow of HCL gas was terminated and thetemperature of the reaction chamber was increased to 1450° C. andallowed to come to equilibrium in about 30 seconds. Silane and propanewere added to the reaction chamber to begin SiC growth. Silane was addedat 200 ppm (Si=200 ppm) and propane was added at 600 ppm (C=1800 ppm) tothe H₂ carrier gas to create a Si/C ratio within the reaction chamber of1:9. A n-type dopant, N₂, was added to the H₂ carrier gas at 40 ppm(N=80 ppm) to begin dopant incorporation into the SiC crystal. A growthrate of 5.5 micrometers/hour was achieved using the improved CVD method.After about one hour, the flow of silane, nitrogen, and propane in theH₂ carrier gas was terminated and the carrier gas was allowed tocontinue to flow in the reaction chamber for about 10 minutes during thecool down of the reaction chamber.

EXAMPLE 2

A method essentially the same as Example 1, wherein a n-type SiC crystalcontaining about 6.0×10¹⁵ cm⁻³ was grown. Silane was added at 200 ppm(Si=200 ppm) and propane was added at 350 ppm (C=1,050 ppm) to the H₂carrier gas to create an Si/C ratio of 1:5.25. A n-type dopant, N₂, wasadded to the H₂ carrier gas at 66 ppm (N=132 ppm) to dope the SiCcrystal.

EXAMPLE 3

A method essentially the same as Example 1, wherein a p-type SiC crystalcontaining about 3.0×10¹⁶ cm⁻³ was grown. Silane was added at 200ppm(Si=200ppm) and propane was added at 350 ppm (C=1,050 ppm) to the H₂carrier gas to create an Si/C ratio of 1:5.25. A p-type dopant,trimethyl aluminum (TMA), was added to the H₂ carrier gas by bubbling H₂into liquid TMA (held at 21° C.) such that 10 sccm (standard cubiccentimeters) was introduced into the reaction chamber to dope the SiCcrystal.

EXAMPLE 4

A method essentially the same as Example 1, wherein a degenerate p-typeSiC crystal for Ohmic contacts containing about 2.0×10¹⁹ cm⁻³ was grown.Silane was added at 200 ppm (Si=200 ppm) and propane was added at 600ppm (C=1800 ppm) to the H₂ carrier gas to create an Si/C ratio of 1:9. Ap-type dopant, TMA, was added to the H₂ carrier gas at 21 sccm to dopethe SiC crystal. About five minutes before cooling the reaction chamber,the silane flow was terminated and H₂, propane, and TMA flows werecontinued.

EXAMPLE 5

A method essentially the same as Example 1, wherein a degenerate n-typeSiC crystal for Ohmic contacts containing about 2.0×10¹⁹ cm⁻³ was grown.Silane was added at 200 ppm (Si=200 ppm) and propane was added at 130ppm (C=390 ppm) to the H₂ carrier gas to create an Si/C ratio of 1:1.95.A n-type dopant, N₂, was added to the H₂ carrier gas at 1153 ppm (N=2306ppm) to dope the SiC crystal. About five minutes before cooling thereaction chamber, the propane flow was terminated and H₂, silane, and N₂flows were continued.

EXAMPLE 6

A method essentially the same as Example 1, wherein a multi-polytype SiCcrystal containing a p/n junction is grown. The 6H--SiC substrate isnucleated with a 3C--SiC crystal. The tilt angle of the substrate was0.2°-0.3°. Crystal growth is begun until the 3C--SiC crystal fans outover the complete growth region. During the initial 3C--SiC growth,degenerate n-type doping, as discussed in Example 5, is carried out. Thedegenerate n-type SiC epilayers are followed by the growth of low dopedn-type epilayer, which are grown substantially the same as epilayersgrown in Example 2. The pn junction is formed by terminating n-typeepilayers and to begin growing degenerate p-type epilayers, as describedin Example 4. The positional distribution of 3C--SiC epilayer mesasversus 6H--SiC epilayer mesas on the substrate was random, and thepercentage of 3C--SiC mesas was roughly 50%. A 2000 Å thick aluminumetch mask defining circular and square diode mesas ranging in area from7×10⁻⁶ cm² to 4×10⁻⁴ cm² was applied and patterned by liftoff. The diodemesas were etched to a depth of approximately 10 μm using reactive ionetching (RIE) in 80% SF₆ :20% O₂ under 300 W rf at a chamber pressure of250 mTorr. The aluminum etch mask was stripped by wet etch, and then acleanup dip in boiling sulfuric acid was performed. The samples were wetoxidized for 6 hours at 1150° C. to form SiO₂ at least 500 Å thick.After the wafers had been patterned for contacts, vias were etched inthe oxide using 6:1 buffered HF solution. Aluminum was then E-beamdeposited and lifted off to complete device fabrication. The 3C diodeexhibits rectification to 200 V reverse bias at ambient temperature (25°C.), which represents a 4-fold improvement in 3C--SiC diode voltagehandling capability. The breakdown is repeatable (i.e., the curve can betaken numerous times with no change in device characteristics) when thecurrent flowing during reverse breakdown is restricted to less than onemilliamp; unlimited current flow results in permanent damage to thediode. On the device tested on the substrate, coronas of microplasmaswere observed exclusively around device boundaries during breakdown,suggesting that reverse failures are occurring at the mesa perimeter andare not due to a bulk mechanism. Once the diode was catastrophicallydamaged by excessive current flow during breakdown, themulti-microplasma corona was replaced by a single microplasma whichpresumably was the point of catastrophic device failure along the mesaedge. FIG. 12 illustrates the forward and reverse current voltagecharacteristics on a logarithmic scale at several temperatures. Althoughthe improvement in reverse leakage naturally depends upon the voltageand temperature selected as a basis of comparison, these 3C--SiC diodesclearly represent at least an order of magnitude improvement in reverseleakage current density over any previously published 3C--SiC pn diode.Because the reverse current is not proportional to the square root ofthe applied voltage, it is surmised that mechanisms other than thermalgeneration are responsible for the reverse leakages. The exponentialregions of the forward characteristics exhibit record-low saturationcurrent densities for CVD-grown 3C pn diodes. However, the change inideality factors with temperature is not well understood at this time.

The method described in the above examples illustrates the controlledand reproducible doping of SiC crystals having a smooth surfacemorphology and a low density of defects. The improved CVD method canpossibly be used to grow and control purity and doping profiles ofcrystals other than SiC having at least two growth sites, even thoughsuch crystals involve a different chemistry, growth parameters, dopants,contaminants, etc. from that of SiC crystals. One type of crystal theimproved method can possibly be used with is GaAs crystals. Thefollowing prophetic examples illustrate the controlled doping of galliumarsenic (Gas) crystals.

EXAMPLE 7

For the method described in this example, a degenerate p-type GaAscrystal containing at least 1.0×10¹⁹ cm⁻³ for Ohmic contacts is to beformed. A GaAs substrate is polished, growth boundaries cut and thesurface etched prior to crystal growth on the substrate. The reactorchamber is brought to the proper crystal growth temperature and thecarrier gas flow is to be brought to equilibrium in the chamber. Galliumand arsenic containing compounds are to be combined with the carrier gasto begin GaAs crystal growth. For the case of a p-type dopant whichpreferentially competes for the Ga site, the gallium concentration is tobe decreased relative to the arsenic concentration, which increases thedemand of Ga growth sites and, therefore, increases the incorporation ofthe p-type into the GaAs crystal. The amount of p-type dopant to beintroduced into the reaction chamber should be the maximum amountallowed which does not adversely affect the GaAs epilayer surfacemorphology. Five minutes before reactor cooling, the gallium compoundflowing into the reactor should be terminated and the flow of the p-typedopant, arsenic compound, and carrier gas continued.

EXAMPLE 8

A method essentially the same as prophetic example 7, wherein adegenerate n-type GaAs crystal containing at least 1.0×10¹⁹ cm⁻³ is tobe grown. For the case of a n-type dopant which preferentially competesfor the As site, the gallium concentration is increased relative to thearsenic concentration to increase the As site demand to increase theincorporation of the n-type dopant into the GaAs crystal. The amount ofa n-type dopant to be introduced into the reaction chamber should be themaximum amount allowed which does not adversely affect the surfacemorphology of the GaAs crystal. Five minutes before reactor cooling, thearsenic compound flowing into the reactor should be terminated and theflow of the n-type dopant, gallium, and carrier gas continued.

EXAMPLE 9

A method essentially the same as prophetic example 7, wherein a p-typeGaAs crystal containing much less than 1.0×10¹⁹ cm⁻³ is to be grown. Thegallium concentration relative to the arsenic concentration is increasedto create a relatively large Ga/As ratio (relative to the Ga/As ratio ofexample 7) in the reaction chamber. An appropriate amount of p-typedopant is then added to the reactor.

EXAMPLE 10

A method essentially the same as prophetic example 8, wherein a n-typeGaAs crystal containing much less than 1.0×10¹⁹ cm⁻³ is to be grown. Thegallium concentration to arsenic concentration is to be decreased in thereactor chamber such that a relatively small Ga/As ratio (relative tothe Ga/As ratio of example 8) exists. An appropriate amount of a n-typedopant is then added to the reactor.

EXAMPLE 11

The SiC epilayers were grown at 1450° C. on commercially availablen-type Si-face SiC substrates polished 3° off-axis from the (0001) planetoward the (1210) plane. The SiC substrates were placed onto aSi--C-coated graphite susceptor and loaded into a water-cooled quartzreactor of an atmospheric pressure CVD system. The SiC epilayers weregrown using silane (3% in H₂) and propane (3% in H₂) with a hydrogencarrier gas, resulting in ˜3 μm/h growth rates, and doped using eithernitrogen (for n-type) or trimethylaluminum (TMA) (for p-type). Theepilayers were characterized using secondary ion mass spectrometry(SIMS) analysis, mercury-probe or p-n diode capacitance voltage (C-V),low-temperature photoluminescence spectroscopy (LTPL), and Hall mobilitymeasurements. Site-competition epitaxy was used by (dopant controltechnique) adjusting the Si/C ratio within the growth reactor toeffectively control the amount of dopant incorporated intosubstitutional SiC crystal lattice sites. These sites, i.e. carbonlattice sites (C sites) or silicon lattice sites (Si sites), located onthe active growth surface of the silicon carbide crystal weremanipulated by site-competition epitaxy. The model for site-competitionepitaxy was based on the principle of competition between nitrogen andcarbon for the C sites and between aluminum and silicon for the Si sitesof the growing silicon carbide epilayer. The concentration of n-type(nitrogen) dopant atoms incorporated into a growing silicon carbideepilayer were decreased by increasing the carbon-source concentration sothat C out competed N for the C sites. Analogously, the amount of p-typedopant (aluminum) incorporated was decreased by increasing thesilicon-source concentration within the growth reactor so that Si outcompeted Al for the Si sites. A series of p-type and n-type doped singlecrystals were grown by varying the propane concentration to effectivelychange the Si/C ratio, while maintaining a constant silane anddopant-source concentration. For the p-type doping of crystals, thepropane concentration was the only parameter varied (400-000 at. ppm) toeffectively change the Si/C ratio within the growth reactor while thesilane (200 at. ppm) and TMA (2 sccm H₂ flow through a constantroom-temperature TMA bubbler) flows were held constant. The p-typedoping revealed the relative increase in silicon concentration (from anincrease in the Si/C ratio) resulted in the exclusion of aluminum from agrowing SiC epilayer. It was found that using a Si/C ratio of 0.1yielded degenerately doped p-type epilayers with net carrierconcentrations of p>1×10¹⁹ cm⁻³. Additional SiC crystals were grownusing a Si/C ratio opf 0.5, which resulted in p-type epilayers with netcarrier concentrations of p=5×10¹⁶ cm⁻³. It was found that as the Si/Cratio was increased from 0.1 to 0.5, the relative amount of Si increasedand out competed the Al for available Si sites of the growing SiClattice, which ultimately resulted in decreased Al dopant incorporation.

During the growth of n-type doped epilayers, the silane (200 ppm) andmolecular nitrogen (90 ppm) concentrations were held constant while thepropane concentration was varied between 133 and 600 ppm. The nitrogendopant profiles were characterized using SIMS and mercury-probe C-V. TheSIMS profile exhibited variations of atomic nitrogen concentrationwithin the grown epilayer from varying only the Si/C ratio (between 0.5and 0.1) during epilayer growth. The net carrier concentration profileobtained from mercury-probe C-V for the same sample correlated well withthe SIMS profile. The results revealed that site-competition epitaxycontrolled the dopant incorporation into electrically active crystalsites of the growing SiC epilayer.

To verify the reproducibility of site-competition epitaxy, a number ofcrystals were grown using the Si/C ratio extremes (0.1 and 0.5) whilemaintaining constant Si (200 at. ppm) and N (100 at. ppm)concentrations. By using a Si/C=0.1 ratio, grown crystals consistentlyresulted in intentionally doped n-type epilayers with net carrierconcentrations of n≈3×10¹⁵ cm⁻³. Further, by using a Si/C=0.5 ratio,grown crystals consistently resulted in n-type epilayers with n≈3×10¹⁷cm⁻³. It is believed that decreased concentration of carbon relative tosilicon (from increasing the Si/C ratio from 0.1 to 0.5) allowed thenitrogen to out compete the carbon for C sites of the growing SiClattice, resulting in an increased atomic nitrogen incorporation.

Very low-doped epilayers were produced by using site-competition epitaxyto exclude the unintentional dopant atoms contained in the growthreactor from incorporating into the grown SiC epilayers. Theseunintentionally doped epilayers, grown using a Si/C=0.1, were examinedusing LTPL spectroscopy to determine the crystalline perfection andrelative concentration of dopant incorporated during growth. The LTPLresults from our lowest unintentionally doped p-type epilayer resultedin an I₇₇ /P₀ =150 and I₇₇ /S₀ -4.7, indicating the most intrinsic SiCreported to date. The dopant concentration in SiC was correlated withthe I₇₇ /P₀ and I₇₇ /S₀ peak-height ratios for low-doped SiC, where P₀and S₀ are the nitrogen donor no phonon line intensities at 3.00 eV andat 2.99 eV, respectively, and I₇₇ is the peak intensity of the SiC77-meV intrinsic phonon replica at 2.947 eV. Using these ratios, a1×10¹⁴ cm⁻³ estimate of unintentional dopant concentration was made byLTPL, which was consistent with the mercury-probe C-V room-temperaturemeasurements of p<5×10¹⁴ cm⁻³. Site-competition epitaxy was used as adopant-exclusion growth technique for the production of low-doped SiCepilayers. Analogously, degenerately doped epilayers were produced byapplying site-competition epitaxy to promote enhanced inclusion ofdopant atoms during SiC epilayer growth. Very thin degenerately dopedp-type and n-type contact layers were formed by ceasing the source flowof Si or C, respectively, during the final 30-40 s of epilayer growthand during substrate cooling. This allowed the dopant atoms toincorporate into the topmost growing p-type or n-type epilayer withoutcompetition from the Si or C atoms, respectively. Degenerate n-typeepilayers were produced by ceasing only the propane flow during theremaining 30-40 s of epilayer growth, while both the silane (200 at.ppm) and nitrogen (300 at. ppm) flows were maintained.

The metal contacts deposited on both the p-type and n-type degenerateepilayers were "ohmic-as-deposited" (i.e. unannealed) for a number ofmetals, which include aluminum, titanium, nickel, and molybdenum.Contact resistivities determined for as-deposited molybdenum andtitanium were calculated linear transmission line method (TLM)! to bepb<5×10⁻⁵ Ωcm² and p_(c) <2×10⁻⁵ Ωcm², respectively, on both n-type andp-type degenerately doped epilayers. At room temperature, thedegenerately doped n-type contact epilayers had measuredlow-electri-field free-carrier concentrations of n=4×10¹⁸ cm⁻³ andyielded mobilities near 60 cm² V s. When incomplete ionization was takeninto account, the n=4×10¹⁸ cm⁻³ low-field free-carrier concentrationtheoretically translates into atomic nitrogen concentrations in excessof N_(D) =2×10¹⁹ cm⁻³. The higher electric fields exerted by the C-Vtechnique (relative to the Hall technique) led to higher measuredfree-carrier concentrations because of field-enhanced ionization thattakes place within the high-field depletion region of a Schottky diode.Relatively lighter-doped n-type epilayers with n=8×10¹⁶ cm⁻³ hadroom--temperature Hall mobilities ranging from 247 to 253 cm⁻² V s andwere found to be less than 10% compensated. Superior SiC electronicdevices were produced by utilizing site-competition epitaxy. Thesedevices exhibited superior electrical performances because the increaseddoping range resulting from utilizing site-competition epitaxy allowedfabrication of SiC devices with much higher blocking voltages and lowerparasitic resistances. Included among the device accomplishments werehigh voltage 6H--SiC diodes (2000 V) and 3C--SiC diodes (300 V) whichexhibited the largest 6H and 3C--SiC blocking voltages reported to date.

EXAMPLE 12

The 6H--SiC epilayers were grown on commercially available n-type (0001)SiC Si-face boule-derived wafers in an atmospheric pressure CVD system.The SiC substrates were precleaned using a standard degreasing solution,followed by immersion in boiling sulfuric acid for 10 min. with a finaldeionized-water rinse and then dried with filtered nitrogen. The cleanedsubstrates were placed onto a SiC-coated graphite susceptor and thenloaded into a water-cooled fused-silica reactor using afused-silica-carrier. The samples were heated via the radio frequency(rf) coupled susceptor which was temperature controlled at 1450° C.using an optical pyrometer. Silane (3% in H₂) and propane (3% in H₂)were used as the sources for SiC epilayer growth, whereas 3% hydrogenchloride gas in hydrogen was used during a 1350° C. in situ 4 min etchjust prior to epilayer growth. All gases were mass flow controlled,including the ultra-pure hydrogen carrier gas which was purified byusing a heated-palladium diffusion cell. The epilayers were doped p-typeby the addition of diborane (100 ppm or 500 ppm B₂ H₆ in H₂) into thereactor during epilayer growth.

Post growth annealing of the crystals was carried out in aCVD-polycrystalline SiC cavity at 1700° C. for 0.5 h in an atmosphericpressure, ultra-pure argon environment. Secondary ion mass spectrometry(SIMS) was performed using a CAMECA IMS-4f double-focusing, magneticsector ion microanalyzer. Cesium bombardment was used for determinationof hydrogen, boron, and nitrogen atomic concentration profiles by usingthe detector in a negative secondary ion detection mode to monitor (H)and the diatomic species B(+C), N(+C), respectively. The aluminum andhigher accuracy boron elemental concentrations were determined by usingoxygen bombardment and a positive secondary ion detection mode.

The amount of boron incorporated into the 6H--SiC epilayers wascontrolled by using site-competition epitaxy. It was first determinedwhether the specific dopant atom occupied either the C-site or theSi-site of the SiC lattice. Boron had been reported to occupy both theSi-site and the C-site, which would preclude effective use ofsite-competition epitaxy for control of boron doping in SiC. However,CVD experimental results indicated that boron preferentially occupiedthe Si-site of the SiC lattice. It was found that the boronincorporation into the SiC epilayer decreases as the silaneconcentration increases (i.e. increasing the Si/C ratio), which wasconsistent with the increased amount of silicon out competing the boronfor available Si-sites during growth of the SiC epilayer. Conversely,boron incorporation was increased by decreasing the Si/C ratio bydecreasing the silane concentration or, alternatively, increasing onlythe propane concentration. The propane concentration was varied foreffective control of the boron doping in which the Si/C ratio wasdecreased stepwise during epilayer growth by successive increases inpropane concentration while maintaining a constant silane and diboraneconcentration. The boron incorporation in the epilayer increased as thepropane concentration was increased due to the relative decrease in theSi/C ratio within the growth reactor. This relative decrease in siliconconcentration enabled the boron atoms to out compete the silicon atomsfor more of the available Si-sites on the surface of the growing SiCepilayer, resulting in increased boron incorporation into the SiCepilayer.

To determine the reproducibility of the doping control, numerousboron-doped SiC epilayers were grown during separate 2 h growthexperiments, each using a constant Si/C ratio ranging from Si/C=0.1-0.5.For selected epilayers, SIMS determined elemental boron concentrationwas compared to the net carrier concentration measured usingmercury-probe C-V. A typical epilayer grown using a Si/C=0.51 had a C-Vmeasured net carrier concentration of p=5×10¹⁵ cm⁻³ as compared to a netcarrier concentration of p=3.5×10¹⁷ cm⁻³ for an epilayer grown using aSi/C=0.11 wherein both crystals grown used an identical diboraneconcentration (1.6 ppm). Secondary ion mass spectroscopy analysisrevealed an elemental boron concentration of B!=6.5×10¹⁶ cm⁻³ for thelower doped (p=5×10¹⁵ cm⁻³) epilayer and a B!=1×10¹⁸ cm⁻³ for the morehighly boron-doped (p=3.5×10¹⁷ cm⁻³) SiC epilayer. The large increase inboron incorporation resulting solely from a change in the Si/C ratioillustrated the strong dependency of the boron incorporation on the Si/Cratio used during epilayer growth, and the preferential B occupancy ofthe Si-site vs the C-site. Three boron-doped SiC epilayers, grown usingidentical Si/C ratios but with different diborane concentrations, weresubsequently examined using low temperature photoluminescencespectroscopy (LTPL). Identical Si/C ratios were used in order toeliminate potential effects of different propane concentrations onhydrogen incorporation, thereby isolating the hydrogen incorporationeffect to only the change in diborane concentration. The resultingspectra reveal that a significant amount of hydrogen was contained ineach of the boron-doped epilayers. The analysis of hydrogen in thecrystal indicated that the hydrogen incorporation was directlyproportional to the amount of boron incorporated into the SiC epilayer.Epilayers containing stepped increases in boron concentration B!, fromonly varying the Si/C ratio, were prepared and subsequently analyzed forhydrogen using SIMS. This was done to confirm that hydrogen and boronincorporation were related, and determine whether the hydrogen wasremoved or if its optical activity was simply altered as a result ofannealing. The increase in hydrogen concentration was observed tocorrespond with the increase in boron concentration within the epilayer.This indicated that the hydrogen incorporation is directly related tothe boron incorporation in the 6H--SiC epilayers. After the 1700° C.anneal for 0.5 h in argon, the sample was again SIMS depth profiled fordetermination of boron and hydrogen concentration. The results indicatedthat the hydrogen had diffused out of the SiC epilayer as a result ofthe 1700° C. anneal. The amount of hydrogen remaining within theboron-doped epilayer was below the hydrogen background concentration(<2×10¹⁷ cm⁻³) in the SIMS instrument. Also, it was noted that boron didnot undergo appreciable solid state diffusion as a result of the 1700°C. anneal, which was evidenced by the continued sharpness of the B!profile. It was found that prior to the anneal, hydrogen was passivatingthe acceptor atoms of the B-doped SiC epilayers. The increase in carrierconcentration, due to the reduction in hydrogen-passivation withpost-anneal, was approximately 4.5×, 3× and 3.25× for the three sampleswith post-anneal net carrier concentrations of p=9×10¹⁵ cm⁻³, p=1.9×10¹⁶cm⁻³, and p=6.5×10¹⁶ cm⁻³, respectively.

It was found that site-competition epitaxy was effective for the controlof B-doping in CVD 6H--SiC(0001) epilayers, with B incorporationinversely proportional to the Si/C ratio used within the reactor duringSiC epilayer growth. Both LPTL and SIMS analysis of the boron-dopedepilayers indicated that the hydrogen concentration contained in theepilayers increased with increasing boron incorporation.

EXAMPLE 13

The 6H--SiC epilayers were grown on commercially available n-type 6H and4H (0001) SiC Si-face and C-face 6H (0001) SiC boule-derived wafer in anatmospheric pressure CVD system. The SiC substrates were precleanedusing a standard degreasing solution, followed by immersion in boilingsulfuric acid for 10 minutes, with a final deionized-water rinse andthen dried with filtered nitrogen. The cleaned substrates were placedonto a SiC-coated graphite susceptor and then loaded into a water-cooledfused-silica reactor using a fused-silica-carrier. The samples wereheated via the RF-coupled susceptor which was temperature controlled at1450° C. using an optical pyrometer. Silane (gas cylinder containing 3%in H₂) and propane (3% in H₂) were used as the sources for SiC epilayergrowth, whereas 90 sccm flow of ultrapure hydrogen chloride gas in a 3sLpm flow of hydrogen was used during a 1350° C. in situ etch just priorto epilayer growth. All gases were mass flow controlled, including theultra-pure hydrogen carrier-gas which was purified by using aheated-palladium diffusion cell. The epilayers were doped n-type by theaddition of phosphine (200 ppm PH₃ in H₂) or nitrogen (0.1% N₂ in H₂)and p-type by the addition of diborane (100 ppm B₂ H₆ in H₂) ortrimethylaluminum (bubbler configuration) into the reactor duringepilayer growth. Secondary ion mass spectrometry (SIMS) was performedusing a CAMECA IMS-4f double-focussing, magnetic sector ionmicroanalyzer. Cesium bombardment was used for determination ofhydrogen, boron, phosphorous, and nitrogen atomic concentration profilesby using the detector in a negative secondary ion detection mode tomonitor H--, P-- and the diatomic species B(+C)--, N(+C)--,respectively. For the 6H-- and 4H--Si(0001) Si-face samples, nitrogendopant incorporation into Si-face epilayers was found to be directlyrelated to the Si/C ratio within the reactor. The Si/C ratio waspurposely varied by changing only the propane flow during the epilayergrowth experiment while both the silane and nitrogen flows weremaintained constant. The SIMS depth profiles indicated that the atomicnitrogen concentration decreased as the propane flow was increasedduring epilayer growth. This was consistent with an increased amount ofcarbon out competing the nitrogen for vacant carbon-sites (C-sites) ofthe SiC crystal lattice on the growing SiC epilayer surface.

Phosphorous (P) dopant incorporation was determined to be inverselyrelated to the Si/C ratio. As the propane flow was decreased, whilemaintaining a constant silane and phosphine flow, the SIMS determined Pincorporation decreased within the Si-face 6H--SiC epilayer. Thesechanges in P incorporation with variation in the Si/C ratio wereconsistent with the P atom occupying the silicon-sites (Si-sites) of theSiC crystal lattice. As the propane flow was increased, the amount ofvacant Si-sites increases which could allow enhanced P incorporationinto the SiC epilayer. The p-type doping on the Si-face was accomplishedusing trimethylaluminum (TMA) for aluminum (Al) and diborane for boron(B) doping. SIMS analysis revealed that the Al dopant incorporation wasinversely related to the Si/C ratio used during epilayer growth on both6H and 4H Si-face substrates. The SIMS determined Al incorporationincreased as the propane flow was increased and also when the silaneflow was decreased within the CVD reactor, which demonstratedcompetition between Si and Al for the Si-sites on the growing Si-face6H--SiC epilayer.

Similarly, boron doping was found to be related to the Si/C ratio. Inaddition to boron incorporation, a significant concentration ofhole-passivating hydrogen was also incorporated during the growth of theB-doped epilayer. SIMS analysis revealed that the hydrogen could beremoved from the epilayer by annealing at 1700° C. in argon for 30minutes which results in a 3× increase in the mercury-probed net holeconcentration for these B-doped epilayers. P-type doping of epilayersgrown on C-face 6H--SiC(0001) samples were found to be similar to thatdetermined for the Si-face substrates. Both Al and B dopantincorporation on the C-face exhibited similar dependence on the Si/Cratio. However, both Al and B incorporation was not as great on theC-face substrate when compared to Si-face substrate that was dopedduring the same growth run. For Al doped C-face epilayers, the SIMSdetermined Al dopant incorporation was approximately 50 times less thanthat of the corresponding Si-face epilayer. Similarly, B-doped C-faceepilayers also exhibited a 50 times lower atomic B concentrationcompared to simultaneously B-doped Si-face epilayers. For B-dopedSi-face epilayers, hydrogen was detected for the B-doped C-faceepilayers, where the variation in the SIMS determined H concentrationprofile paralleled that of the SIMS determined, stepped B-concentrationprofile.

It was found that Si-face site-competition for 6H--SiC(0001) was similarto that of the 4H--SiC(0001) substrates. The site-competition effect forAl doping was consistent with Al substantially occupying the Si-sites ofthe SiC lattice for epilayers grown on both Si-face and C-face6H--SiC(0001) substrates. A similar site-competition effect was obtainedin which B was believed to primarily occupy the Si-sites for B doping onboth faces. For n-type doping with phosphorus, the site-competitioneffect was most consistent with the P atom mainly occupying theSi-sites. This effect is believed to be best explained by consideringthat the lattice site occupied by each dopant atom as compared to thesize of the Si or C atom. Specifically, the non-polar covalent radii forthe elements of interest are: Si(1.17A); C(0.77A); Al(1.26A); B(0.82A);P(1.10A) and N(0.74A). Using atomic size as a first approximation (i.e.by neglecting chemical bonding arguments), both Al and P shouldsubstitute for Si and not for C in the SiC lattice, which is consistentwith the experimental results. The more closely matched atomic size of Bto C suggests that B should substitute mainly for C in the C-sites,which is in direct conflict with experimental evidence. However, aconsiderable amount of H is simultaneously incorporated into the growingB doped epilayer and therefore a "B--H" complex explains these results.A "B--H" species has a size (1.10A) more closely matched with that ofSi(1.17A) and therefore should occupy a Si-site, which is in agreementwith the experimental results. This B--H complex reveals that theSi-site substitution is favored because of the larger effective size ofthe proposed "B--H" species.

The N doping on the Si-face was consistent with the site-competitionmechanism, in which the relatively small N atoms compete with C atomsfor available C-sites. It was also found that the small size of N alsoenabled N to compete to a less extent for the Si-sites. This effect ismost pronounced when the SiC crystal is grown on the C-face of asubstrate. The results of using site-competition on the C-face wereconsistent with the N-dopant atoms preferentially substituting into theSi-sites, with a significant contribution from N incorporation into theC-sites.

Site-competition epitaxy was successfully used for p-type and n-typedopant control on both the Si-face and C-face of 6H--SiC(0001) off-axissubstrates. The doping results for the C-face indicated that thesite-competition effect was highly dependent on the SiC polarity but notsignificantly affected by the polytype. The results for both the 6H and4H SiC Si-face samples were similar whereas the 6H SiC C-face samplesyielded significantly different results, specifically for N-dopedepilayers. In comparing the results from simultaneous growth of N-dopedepilayers on the C-face and Si-face, that the site-competition mechanismwas believed to be dependent on the surface structure of the SiCsubstrate and relatively less dependent on gas phase interactions. Forp-type doping, the Al and B dopant incorporation increased as the Si/Cratio was decreased on both polar faces. Similarly, the P dopantincorporation also increased on both the Si-face and C-face as the Si/Cratio was intentionally decreased. This was consistent with the P, Al,and B competing with Si for the available Si-sites. The site preferencefor each dopant was found to be dependent on its atomic size relative tothe size of Si(Si-site) and C(C-site). Here, the P and Al were clearlytoo large for incorporation into the C-site without changing the latticespacing. Because a significant amount of H was incorporated with B, therelative size of B to be of a "B--H" complex was considered. For thissituation, the effective size of the "B--H" complex was also too largeand would not easily fit into the C-sites without significantly alteringthe SiC lattice dimensions. In contrast, the relatively small size ofthe N dopant atom allows incorporation into either of the lattice sites.Therefore, N doping on the Si-face results in N incorporation into boththe Si-site and the C-site, with N mainly incorporating into the C-site.In contrast, for N doping on the C-face, N mainly incorporates into theSi-sites with relatively less incorporation into the C-sites. Therefore,the site-competition mechanism for N dopant incorporation into bothSi-sites and C-sites, is apparently dependent upon the growth surfacepolarity.

EXAMPLE 14

The 3C--SiC and 6H--SiC epilayers were grown at 1450° C. on commerciallyavailable 6H(0001) silicon-face SiC substrates with tilt angles rangingfrom 0.2° to 0.6° and 3° to 4°, respectively. The SiC substrates wereplaced onto a SiC-coated graphite susceptor and loaded into awater-cooled quartz reactor CVD system. The grown SiC epilayers wereexamples using mercury-probe or pn diode CV to determine active dopingconcentrations. The Si--C ratio within the growth reactor had a stronginfluence on intentional and unintentional dopant incorporation of thegrown 6H, 3D, and 4H SiC(0001) epilayers. Specifically, the activen-type (nitrogen) carrier concentration was found to be directlyproportional to the Si/C ratio, whereas, the active p-type (aluminum)concentration was found to be inversely proportional to the Si--C ratiofor epilayer growth on the SiC(0001) basal plane. As the Si/C ratiowithin the growth reactor was decreased, the active nitrogenconcentration in the grown SiC epilayer decreased. When the Si/C ratiowas decreased further to Si/C≈0.1, the unintentionally doped epilayerschanged from n-type to p-type. As a result, we have obtained both p-typeand n-type epilayers with room temperature carrier concentrations of1×10¹⁴ cm⁻³, as measured by both mercury-probed CV and low temperaturephotoluminescence (LTPL). Previously, the unintentionally dopedepilayers produced in our lab were exclusively n-type with the lowestnet carrier concentrations typically limited to about n=2-3×10¹⁶ cm⁻³.

Dopants in SiC are believed to occupy specific sites, specificallynitrogen primarily occupies the carbon site on the Si-face whilealuminum occupies the silicon site of the SiC lattice on both sides ofthe crystal face. The relative increase in carbon concentration "outcompetes" the nitrogen for the C-sites of the growing SiC lattice. Theanalogous situation exists for an increased Si/C ratio, in which therelative increase in silicon concentration "out competes" the Al for theSi-sites of the growing SiC lattice. 100 ppm of N (50 ppm N₂) wasintroduced into the growth reactor and then attempted to exclude it fromthe growing epilayer by only decreasing the Si/C ratio from abutSi/C=0.44 to Si/C=0.1. The Si/C of 0.1 results in consistently producingan intentionally doped n-type SiC epilayer with a net carrierconcentration of n=1×10¹⁵ cm⁻³. In contrast, growth using the moretypical Si/C=0.44 with 100 ppm of N results in n-type epilayers ofn=1-2×10¹⁷ cm⁻³. For epilayer growth using the Si/C ratio of 0.1, theincreased amount of carbon is believed to have out competed the nitrogenfor the C-sites of the growing SiC lattice. P-type epilayers were alsoproduced using this method. For these series of experiments a flow ofTMA was introduced into the reactor during epilayer growth. Theresulting p-type epilayer grown with a Si/C ratio of 0.44 was measuredto be 5×10¹⁶ cm⁻³, while the epilayer grown using a Si/C ratio of 0.1yielded a degenerately doped p-type epilayer with an estimated netcarrier concentration of 1×10¹⁹ cm⁻³. As the Si/C ratio was decreasedform 0.44 to 0.1 the relative amounts of Si competing with the Al forthe Si-sites of the SiC lattice also decreased, which resulted in anincreased Al incorporation.

Site-competition epitaxy was also successfully used to obtain veryabrupt changes in dopant concentrations in SiC epilayers. Inconventional CVD systems, the abruptness of the dopant profile islimited by the purging of the dopant-source from the growth reactor. Oneadvantage of epilayer growth using site-competition epitaxy is that moreabrupt dopant profiles can be obtained by excluding the remainingunwanted dopant by changing the Si/C ratio along with the dopant-sourcegas. Conversely, very abrupt, enhanced dopant incorporation can beaccomplished for production of highly degenerately doped epilayers. Onehighly useful example of this was the formation of very thindegenerately doped p-type and n-type contact layers by ceasing thesource-flow of Si or C, respectively, during the last minutes ofepilayer growth. Subsequently deposited metal contacts were "ohmic asdeposited" for a wide variety of metals on both p-type and n-typeepilayers. The contact resisivities were determined for as depositedmolybdenum on both n-type and p-type degenerately doped epilayers werecalculated to be p_(c) <5E-5Ωcm² using the linear TLM method.Preliminary Hall measurements on intentionally doped 6H--SiC epilayersamples were carried out. N-type epilayers with n=8×10¹⁶ cm⁻³ were foundto be less than 10% compensated with room temperature Hall mobilitiesranging form 247 to 253 cm² /V-s. At room temperature, the heavily dopedn-type contact epilayers yielded mobilities near 60 cm² /V-s andmeasured low-electric-field free carrier concentrations of n=4×10¹⁸cm⁻³. When incomplete ionization was taken into account, the 4×10¹⁸ cm⁻³low-field free carrier concentration theoretically translated intoatomic nitrogen concentrations in excess of 2×10¹⁹ cm⁻³. The"ohmic-as-deposited" state of the contacts precluded conventional C-Vprofiling of the heavily doped epilayers. For the lowest doped n-typeand p-type epilayers, low temperature photoluminescence (LTPL) wasemployed to determine crystalline perfection and relative dopantincorporation. The LTPL results from our lowest unintentionally dopedp-type epilayer resulted in a I₇₇ /P₀ =150 and an I₇₇ /S₀ =4.7. Theestimate of unintentional dopant concentration by LTPL (1×10¹⁴ cm⁻³) wasconsistent with the mercury-probe CV room temperature measurements of p5×10¹⁴ cm⁻³.

The novel growth method, based on the use of appropriate Si/C ratiosduring epilayer growth to affect control over dopant concentration, wasdemonstrated for SiC(0001) basal plane 6H and 3C--SiC. The carrierconcentration change as the Si/C ratio was varied from about 0.1 to 0.8.the donor carrier concentration in the grown epilayer was proportionalto the Si/C ratio whereas the acceptor carrier concentration wasinversely proportional to the Si/C ratio. The resulting surfacemorphologies and crystal quality were excellent for all Si/C ratiosbetween 0.1 to 0.8.

OVERVIEW OF SOME APPLICATIONS

The electrical device applications of the invention center around theadvantageous use of the wider doping ranges that this new growthtechnique makes possible. Since the invention allows for theincorporation of higher concentrations of dopant into compoundsemiconductor crystals than was previously possible, any electricaldevice structure whose performance is affected by the degree ofdegenerate doping achieved in any region of the device will beinfluenced by the present invention.

The improved CVD method significantly reduces contact resistances at theOhmic contacts. Performances of many semiconductor device structures,such as transistors, are dependant on contact resistance. Contactresistance arises due to the physical interface where the metal (whichcarries electrical signals and power to and from the semiconductordevice) contacts (or connects to) a semiconductor surface. Theperformance of most semiconductor electrical devices are maximized byminimizing the contact resistances. It is a limiting factor in theelectrical capability of many devices, such as transistors, which inturn limit the capabilities of electrical circuits and systems madeusing those devices. It is widely recognized that the resistance ofcontacts to 6H--SiC or 3C--SiC is the major limiting factor in theelectrical capability of many 6H--SiC or 3C--SiC devices respectively.The contact resistances for semiconductor devices are minimized bemaximizing the doping in the semiconductor where it makes physicalcontact with metal Ohmic contact. Because the improved CVD methodenables device quality (i.e. high quality) epilayers of compoundsemiconductors (like SiC for example) to be grown with higher dopingconcentrations, contact resistances to these epilayers (and thereforeproperly designed electrical devices fabricated in these epilayers) canbe reduced thereby increasing the performance capabilities of manycompound semiconductor based electrical devices and circuits.

The improved CVD method also reduces bulk semiconductor resistances inthe grown crystals. Another source of resistance that can affectsemiconductor electrical device performance is the resistance of thebulk semiconductor itself. In almost all semiconductor devices, currentflows through undepleted regions of the bulk semiconductor. Theresistance associated with charge flow through these regions is oftensignificant and can sometimes be an undesired factor which limits deviceperformance. The resistances associated with these regions can beminimized by maximizing the doping of the semiconductor in theseregions. The maximization of doping in these regions throughconcentration levels previously unattainable is again facilitated bythis new inventive method.

The present invention can also be used to form Delta-doped semiconductordevices. These structures are based upon very thin, very heavily dopedsemiconductor layers. The electrical performance of Delta-doped devicesis contingent upon the ability to incorporate a maximum amount of dopantinto the thinnest possible layer. The ability to better control thelevel of degenerate doping of compound semiconductors, which is madepossible by the improved CVD method, should significantly enhance theperformance characteristics obtainable in Delta-doped device structures.

The improved CVD method further can be used to develop semiconductordevices with smaller depletion widths. The depletion widths found withinsemiconductor devices at various junctions (e.g. pn junctions, metalsemiconductor junctions, hetero-junctions) are largely determined by thesemiconductor doping at or around the junction. It is well known thatthe depletion width decreases as the doping concentration increases.There are many semiconductor devices whose performance and/or functionrelies on heavy doping to obtain narrow depletion widths and theassociated physical effects. Zener diodes and Esaki diodes rely oncarrier tunneling through a very narrow depletion region made possibleby heavy doping on both sides of a pn junction. Low leakage diodes anddiode junction charge storage capacitors are based on narrow depletionwidths, so as to minimize generation current and maximize charge storagedensity. Because the improved CVD method enables heavier doping andresults in smaller depletion widths, the improved CVD method has thepotential to improve the performance of these devices by narrowing thedevice depletion widths.

The improved CVD method can also be used to develop semiconductordevices having slightly narrower bandgaps. By increasing the degeneratedoping in a semiconductor, a physical phenomenon known as band tailingor bandgap narrowing takes place. The improved CVD process can be usedto control the degree of degenerate doping in semiconductors used indevices in which bandgap narrowing enhances device performance, such aslight-emitting diodes (LEDs) and bipolar transistors.

In addition to the controlled growth of high quality compoundsemiconductor epilayers with higher doping concentration, the improvedCVD process also enables the controlled growth of high quality compoundsemiconductor epilayers with lower dopant concentrations than waspreviously possible. Any compound semiconductor electrical device whoseperformance stands to gain from the lowered doping concentrations thatthis improved CVD method enables will benefit. Such devices which willbenefit include fundamental semiconductor junctions. Almost allsemiconductor devices and circuits contain fundamental semiconductorjunctions. Most transistors composed of these fundamental junctions arearranged in various configurations appropriate to accomplish the desiredelectrical functions (e.g. amplification, switching). Such devices willbe greatly improved by the improved fundamental junction doping of thepresent invention. One fundamental building block of semiconductorelectrical device technology is the intimate junction of p-type materialto n-type material, more commonly referred to as p-n junction. It iswell known to those skilled in the art that the electricalcharacteristics (such as junction capacitance, junction breakdownvoltage, junction leakage current) of the p-n junction are governed bythe physical characteristics (i.e. dopant density, defect density) ofthe lighter-doped side of the junction. The ability to produce a morelightly doped compound semiconductor, while obtaining the desiredpolarity (n-type or p-type) via this improved CVD method, dramaticallybroadens the range of electrical operating characteristics possible indevices that incorporate pn junctions. It is well known that thebreakdown voltage of a p-n junction diode increases with reductions inthe doping of the lighter doped side of the junction. Therefore, afactor of 10 or greater improvement (i.e. reduction) in dopingconcentrations has been demonstrated by the improved CVD method, whenapplied to 6H--SiC crystals, could improve 6H--SiC p-n step junctiondiode blocking voltages from their current demonstrated maximum ofslightly more than 1,100 volts to somewhere around 10,000 volts. In asimilar manner, the doping of the semiconductor also plays a major rolein the electrical characteristics of another fundamental junction, themetal semiconductor junction. Similar improvements with similarramifications can be expected for properly designed rectifying metal-SiCSchottky diode junctions. The improved CVD method can also be used inthe design of transistors and circuits to permit higher circuitoperating voltages.

The improved CVD method can also be used to make semiconductors withreduced junction capacitance. It is well known that diode junctioncapacitance decreases as doping decreases. The lighter doping, madepossible by the improved CVD method, can also lead to the reduction inthe depletion capacitance of compound semiconductor junctions, whichwill reduce the parasitic capacitances that can often limit device andcircuit performance. One of these factors that limits the switchingspeed and high frequency performance of planar Field Effect Transistors(FET's) is the parasitic capacitance formed by the drain-to-substratep-n diode junction. The decreased doping obtainable by this improved CVDmethod can decrease this junction capacitance and possibly enablecompound semiconductor FET's to operate with higher frequencies andswitching speeds than previously possible.

This improved CVD method can further be used to form semiconductordevices with high internal resistivity. It is known that resistivityincreases as shallow ionization energy dopant concentration decreases.In many semiconductor device and integrated circuit applications, it isadvantageous to build devices in high resistivity (low dopantconcentration) material. Such material often provides isolation betweenadjacent devices on the same chip, and the degree of isolationdetermines the spacing between adjacent devices within a chip, which isa big concern on integrated circuits with hundreds of thousands oftransistors. Since the degree of isolation is in part a function of thesemiconductor purity (i.e. doping density and defect density) betweenthe individual devices, the improved CVD method may prove useful inadvantageously shrinking device spacing to allow for more devices on asingle chip. Clearly, the method is useful for any device or circuitsituation where higher-resistivity compound semiconductor epilayers aredesired.

The improved CVD method can also be used to grow crystals used insemiconductor devices in which the crystals have increased carriermobilities. It is known that carrier mobilities increase as doping anddefect concentrations decrease. This is due to the fact that there isless potential perturbations (caused by impurities and defects) for acarrier (electron or hole) to "plow into" as it is moving through thecrystal lattice under the influence of an electric field. The fact thatthe improved growth technique enables the growth of purer material withlower unintentional dopant incorporation should increase the carriermobility in the compound semiconductors. It is well known to thoseskilled in the art that carrier mobilities directly affect theperformance of many devices and circuits, and that increased carriermobilities are attractive because they enhance the performance of mostsemiconductor devices, especially transistors. In most transistors,crucial electrical performance factors (switching speed, maximumoperating frequency, current carrying capability, gain) are generallyenhanced by increased carrier mobilities.

The improved CVD method can be used to grow crystals having increasedcarrier lifetimes. Carrier lifetime is another physical property ofcompound semiconductors that can be enhanced by growing purer crystalswith lighter background defect concentrations. It is known to thoseskilled in the art that the bulk carrier lifetime increases as thecrystal defect concentration decreases. An improvement in this propertycould lead to smaller recombination/generation rates within a deviceresulting in improved performance for certain types of devices, such asminority carrier based devices. Minority carrier based devices, such asbipolar transistors and solar cells for example, could benefit fromlonger life-times.

Possible industrial uses of SiC crystals grown using the improved CVDmethod includes semiconductor devices and sensors for use in hightemperature environments, such as advanced turbine engines, space powersystems, deep-well drilling, advanced automobile engines, etc.;semiconductor devices and sensors for use in high radiationenvironments, such as found near nuclear reactors; semiconductor devicesfor power electronics applications, such as will be required for powerconditioning electronics for electric vehicles, space power systems, andfor electrical actuators on advanced aircraft; semiconductor devices forhigh frequency applications, such as found in communication satellites,high speed computers, and microwave power transistors used in radarsystems; pressure transducer diaphragm material for high temperatureand/or corrosive environments; and, light-emitting diodes (LEDs) for usein high temperature environments.

As stated above, a significant improvement in producing doped crystalsis the controlling of the contaminant content and/or dopantconcentration of a crystal during crystal growth by varying theconcentration of crystal source components during crystal growth.Several techniques which have been previously used to dope crystalsformed by chemical vapor deposition (CVD) have significant drawbacks anddo not form a crystal which can be repeatedly duplicated and has acontrolled amount of dopant as the crystals formed by the processdisclosed and claimed in the present invention. One of the majorlimitations to progress in silicon carbide (SiC) crystal growth is thelimited, reproducible doping range available. The prior art doping ofSiC epilayers with net carrier concentrations of less than 3×10¹⁶ cm⁻³or greater than 5×10¹⁸ cm⁻³ was not possible, until now, in areproducible, controllable manner. The lower bound on the doping range(3×10¹⁶ cm⁻³) limited high voltage electronic applications. The upperbound on doping concentration (5×10¹⁸ cm⁻³) necessitated the need formetal contacts to be annealed using specific high temperature processes.In contrast to prior art teachings, the present invention provides for agreatly expanded, reproducible doping range from 1×10¹⁴ cm⁻³ or less togreater than 1×10¹⁹ cm⁻³. Very low doped epilayers are needed forproduction of electronics which can withstand much higher voltages thanwere previously available. Very high doped epilayers allow for metalcontacts which do not require the prior art process of high temperatureannealing.

The significant improvement in controlling the amount of contaminantand/or dopant in a crystal grown in a vapor deposition process wasachieved by manipulating the site competition for a particularcontaminant and/or dopant (site-competition epitaxy). Many types ofcrystals can be grown by the CVD method such as a silicon carbidecrystal. The growth of silicon carbide crystals is accomplished by usinga source of silicon (Si) and a source of carbon (C), where the Si atomsoccupy only Si-sites and C atoms occupy only the C-sites of the siliconcarbide (SiC) crystal lattice. The Si sites are also calledSi-growth-sites, just as the C-sites are also referred to asC-growth-sites. When a silicon carbide crystal is doped, a dopant atomsuch as phosphorous, nitrogen, aluminum, boron, etc. substitutionallyoccupies a Si-site or a C-site in the SiC crystal lattice. It has beenfound that the amount of dopant which occupies a particular Si-site or aC-site in the SiC crystal lattice can be controlled by manipulating theSi/C concentration ratio within the CVD reaction chamber. In particular,it was found that specific p-type dopant incorporation can be increasedor decreased by varying the Si/C concentration ratio while maintaining aconstant dopant concentration within the CVD reaction chamber. Anincrease in the Si/C concentration ratio allows the Si to compete with adopant atom for the Si-site while a decrease in the Si/C concentrationratio allows the C to compete with a dopant atom for the C-site. Bymanipulating the Si/C concentration ratio, the amount of dopantincorporated into a growing SiC crystal can be accurately controlled byadjusting the Si/C concentration ratio. The manipulation of the Si/Cconcentration ratio directly affects the site competition at both the Sisite and the C site. It was also found that site competition epitaxy canbe used for both sides of a substrate (i.e. Si-face and C-face for Si(substrate)). For SiC crystals, site competition epitaxy has beensuccessfully used for various types of crystal lattices (i.e. GH, 4H and3C). By using the doping method disclosed and claimed in the presentinvention, crystals grown by CVD process having a particular dopantconcentration can be accurately duplicated in subsequently growncrystals. Prior methods of crystals grown by CVD process failed toreliably produce duplicative crystals having a uniform dopantconcentration. By use of site-competition epitaxy, multiple numbers ofcrystals having an almost identical dopant concentration can be producedhaving a concentration which is significantly lower (when lower isdesired) and significantly higher (when higher is desired) than dopantconcentrations previously obtained by prior methods. By properlytreating the substrate onto which the crystal is to be grown, the growncrystal will be of a higher quality with much fewer defects. Suchpre-treatment includes etching, polishing and forming grooves on thesubstrate to form a desired surface for crystal growth. The substratemay be intentionally nucleated so as to manipulate the growth directionof the crystal and to facilitate in the start of crystal growth duringthe CVD process. The substrate may also be angled to facilitate acertain type of crystal growth.

Further modification of the invention is evident. For instance, thecrystal element ratio can be selected to effectively exclude aparticular dopant atom from a growing crystal. This can be accomplishedwithin a broad doping range. For example, the exclusion ofunintentionally present Al (i.e. in the reactor) from incorporation inton-type epilayers to improve MOS device characteristics were achievedusing site-competition epitaxy techniques. The prior art is absent anyteachings for doping crystals (i.e. SiC) to exclude compensatingdopant-atoms (e.g. Al in N-doped layers) from incorporation into agrowing crystal. A 5E16 cm⁻³ nitrogen doped (n-type) epilayer can growusing a specific flow of nitrogen (n-type dopant) for each reactorduring epilayer growth. The specific nitrogen flow varies only slightlyto allow for variations between growth reactors and other equipmentdifferences. Once determined, this "recipe" then becomes the onlyspecific method for a particular reactor to produce a specific dopantconcentration exhibiting a specific electrical character. Theincorporation of unintentionally present Al into the crystal adverselycompromises the electrical characteristics of the crystal. Until thepresent invention, these limitations could not be overcome. It has beendiscovered that doping such as with nitrogen to obtain a n=5E16 dopinglevel can be accomplished using a spectrum of nitrogen flows, eachcoupled with the appropriate Si/C ratio. Therefore, the same n-typedoped epilayer (i.e. 5E16) can be accomplished by using a wide range ofspecific Si/C ratios with a range of nitrogen flows. Thus, a method forthe exclusion of counter-dopant atoms which are typicallyunintentionally present (residual) in the growth reactor can beachieved. Therefore, a specific Si/C ratio can be chosen which excludesa specific unwanted counter-dopant (aluminum as a p-type counter-dopant)by using a large Si/C ratio while achieving the needed n=5E16 n-typedoping (i.e. nitrogen doped) level.

This is important if a p-type counter-dopant is unintentionally present(as an impurity) and needs to be excluded from the intentionally n-typedoped epilayer. Specifically, aluminum has been reported to adverselyaffect MOS (metal-oxide-semiconductor) device performance in Si--C-baseddevices by causing formation of defects suspected in the oxide withinthe SiO2/SiC MOS-structure. N=5E16 n-type epilayers produced using arelatively large Si/C ratio (to exclude residual Al) improved MOS deviceperformance whereas otherwise identical n=5E16 n-type epilayers producedusing a relatively smaller SiC ratio resulted in decreased MOS-deviceperformance. Also, samples which were grown using a large SiC ratiocontained less residual (unintentional) Al whereas large amounts of(unintentional) Al were incorporated into samples grown with relativelysmall Si/C ratios.

It will be appreciated that multiple "recipes" for crystals such as SiCcrystals can be prepared for various crystal doping ranges. For example,it has been established that single CVD grown SiC crystals can be grownand doped using a wide range of Si/C ratios and corresponding dopantflow rates. The concept of using a multiple of Si/C ratios to produce acrystal with a specific amount of dopant incorporation is contrary tothe teachings in the prior art of selecting a single optimum Si/C ratioand exclusively controlling the dopant concentration by varying the flowrate of dopant into the growth chamber. These "recipes" can be stored ina computer database or the like for dopant control of the grown crystal.As described above, if a SiC crystal is to be grown with n-type dopantincorporation of a specific concentration, a recipe can be selected togrow such a crystal. Further, if a very pure n-type SiC crystal isneeded, a specific recipe is selected which adjusts the selected Si/Cratio and n-type dopant flow rate such that the proper amount of n-typedopant is incorporated into the Si/C crystal but contaminants such as Alare excluded from being incorporated in the SiC crystal. Alternatively,if a very pure p-type SiC crystal is required, a recipe is selected fora particular Si/C ratio and p-type dopant flow rate such that the properamount of p-type dopant is incorporated into the Si/C crystal butcontaminants such as N are excluded from being incorporated in the SiCcrystal. As can be appreciated, if a certain amount of p-type and n-typedopant is needed in the grown crystal, the crystal element ratios can beset accordingly. The significant discovery that the amount of dopantincorporation and/or contaminant exclusion is dependent on the crystalelement ratio is a significant improvement of regulating crystal doping.Not until the invention has there been a method to exclude contaminantsfrom growing in CVD grown crystals. The control method is especiallyimportant for low dopant concentration wherein contaminants incorporatedinto the crystal have a much greater adverse affect to the electricalproperties of the crystal.

Another embodiment to the invention is the use of dopant to excludecontaminant incorporation into the CVD grown crystal. For example, arecipe for a crystal was selected to intentionally out-compete dopantatom-A (unintentionally present) by intentionally introducing dopantatom-B found to decrease incorporation of dopant-A into the growingsingle crystal epilayer. This example has advantages in applications inwhich certain dopant atoms (present even in background concentrations)adversely affect the performance of a final device. For instance,aluminum and boron are typically trace impurity dopants found inintentionally-nitrogen doped n-type epilayers. In addition to theunwanted electrical-compensation effect of aluminum and boron onnitrogen, aluminum has been related to poor MOS device performance. Asolution to this problem of unintentional aluminum incorporation (whichleads to poor MOS devices) was the following:

During nitrogen-doped n-type epilayer growth, the use of an additionaln-type dopant atom (such as phosphorus) which competes for the samelattice sites as the unwanted p-type dopant atom (i.e. Al) wassimultaneously introduced into the reactor with the nitrogen dopant(which incorporates into a site different from the P and Al).Specifically, phosphorus was simultaneously introduced into the reactorduring crystal growth to exclude any residual aluminum fromincorporation into the n-type (mainly accomplished using nitrogen as then-type dopant) epilayer. The exclusion of Al was successful because bothphosphorus and aluminum compete for the Si-lattice sites in SiC. Theincorporation of phosphorus did not adversely affect the crystal becausephosphorus is also a n-type dopant atom. The mixture of phosphorus andnitrogen as intentional n-type dopants provided a method to exclude theunwanted Al as a residual, compensating p-type dopant. In addition tophosphorus, fluoride, chlorine and sulfur can also be used to excludealuminum from Si-sites.

As will be appreciated, other p-type dopants such as boron, sodium,iron, gallium, titanium and vanadium which occupy Si-sites will beexcluded during n-type doping by use of n-type dopants which occupySi-sites such as nitrogen, fluorine, chlorine and/or sulfur.Alternatively, n-type dopants such as nitrogen, fluorine, phosphorus andsulfur will be excluded from Si-sites during p-type doping by use ofsodium, iron, aluminum, gallium, titanium and vanadium. By use of thismethod, a very pure p-type doped crystal will be grown by using a p-typedopant which occupies C-sites such as boron, selecting a Si/C ratiowhich excludes unintentional n-type dopant incorporation into Si sites,and introducing a p-type dopant which competes for Si-sites (i.e. Na,Fe, Al, Ga, Ti, V) to further exclude n-type dopant incorporation intothe crystal.

A brief list of applications which utilize this method are as follows:

1. Use of P (phosphorus) to dope n-type to exclude:

a. Al (p-type) from Si-sties (i.e. 1) reduce the electrical compensationof nitrogen doped n-type epilayers; 2) to further hasten the change inpolarity during growth of crystal epilayers from p-type to n-type).

b. N from Si-sites; especially on C-face samples (also on Si-face and Aface samples) because N incorporation from the C-sites could bepreferred over N incorporation into the Si-sites.

c. N from Si-sites and simultaneously using relatively highSi-source/C-source ratios for better exclusion of N from incorporationinto Si-sites.

d. B (Boron) from Si-sites to force B mainly into C-sites; 1) to excludeH from incorporation into crystal (H is incorporated only when B is inthe Si-site but not when B is in the C-site).

e. Ga (gallium), Ti (titanium); V (vanadium), Na (Sodium) from Si-sites.

2. Use of B (boron) to dope p-type to exclude:

a. Al from Si-sites, thereby excluding Al (intentional and/orunintentional) from incorporation into crystal.

b. N from Si-sites, especially on the C-face but also some on theSi-face, thereby forcing N incorporation into C-sites while excluding Nfrom Si-sites.

3. Use of O (oxygen) to dope n-type to exclude:

a. B from C-sites.

b. N from C-sites because N incorporation into Si-sites could bepreferred over N incorporation into C-sites.

4. Use of P (phosphorous) to dope n-type to exclude:

a. Na from Si-sites.

b. S from Si-sites.

Another modification of the invention is the selection of a particularcrystal face in combination with site-competition epitaxy. Inparticular, dopant incorporation rates have been found to be dependenton the particular face of the crystal. For instance, nitrogen primarilyincorporates into Si-sites on the C-face of the SiC substrate. However,nitrogen primarily incorporates into C-sites on the Si-face and A-faceof a SiC substrate. Consequently, nitrogen can be used to exclude p-typedopants such as aluminum from Si-sites when the crystal is grown on theC-face of a substrate. Conversely, nitrogen can be excluded from ap-type crystal by growing the crystal on the C-face and adding Al, Na,Fe, Ga, Ti and/or V in the growth chamber to exclude N from the Si-site.

The invention has been described with reference to a preferredembodiment and alternates thereof. It is believed that manymodifications and alterations to the embodiment as discussed herein willreadily suggest themselves to those skilled in the art upon reading andunderstanding the detailed description of the invention. It is intendedto include all such modifications and alterations insofar as they comewithin the scope of the present invention.

I claim:
 1. A method of regulating an amount of at least one non-crystalelement deposited in a given growth area of a single crystal formed fromat least two crystal elements as said crystal is grown in a growthchamber, said crystal elements comprising a first crystal element and asecond crystal element, said crystal having at least two crystal growingsites wherein said first crystal element is deposited in a first growthsite and said second crystal element is deposited in a second elementgrowth site, said at least one non-crystal element being competitive forsaid first site, said method comprising the steps of:a. selecting anamount of said non-crystal element to be deposited in said given growtharea; b. selecting a ratio of said first crystal element to said secondcrystal element, said ratio being dependent on said selected amount ofsaid non-crystal element to be deposited in said given growth area; c.flowing a controlled amount of at least one gaseous crystal elementcompound through said chamber to grow said crystal, said controlledamount corresponding to said selected ratio of said first crystalelement to said second crystal element in said growth chamber, and eachof said element compound including at least one of said crystal element;d. selecting a new amount of said non-crystal element to be deposited insaid given growth area; e. selecting a new ratio of said first crystalelement to said second crystal element to obtain said new amount of saidnon-crystal element to be deposited in said given growth area; and, f.changing said ratio of said crystal elements in said growth chamberduring the growing of said crystal to regulate said amount of said atleast one non-crystal element deposited in said competitive crystalgrowth site at said growth area to obtain said new amount of saidnon-crystal element being deposited in said given growth area.
 2. Amethod as defined in claim 1, including a step of changing the amount ofsaid at least one non-crystal element being deposited in said firstgrowth site by changing said ratio of said second crystal element tosaid first crystal element.
 3. A method as defined in claim 2, whereinsaid at least one non-crystal element is a dopant.
 4. A method asdefined in claim 3, including a step of pre-treating said growth area toremove sites for heterogeneous nucleation during said crystal growth. 5.A method as defined in claim 4, wherein said pre-treating includespregrowth etching of said growth area.
 6. A method as defined in claim4, wherein said pre-treating includes polishing said growth area.
 7. Amethod as defined in claim 4, wherein said SiC crystal is grown on thegrowth area of a substrate.
 8. A method as defined in claim 7, includinga step of nucleating on said substrate to grow a particular crystalpolytype.
 9. A method as defined in claim 8, wherein said nucleating ofsaid substrate is the introduction of an impurity at said site tostimulate the growth of a desired crystal structure.
 10. A method asdefined in claim 9, wherein said nucleating of said substrate ispositioned in a corner of said growth area.
 11. A method as defined inclaim 7, wherein said growth area of said substrate has a non-zeroinclination relative to a basal plane.
 12. A method as defined in claim11, wherein said non-zero inclination is greater than about sevendegrees.
 13. A method as defined in claim 11, including a step ofdividing said growth area into growth regions defined by boundaries byintroducing grooves along the boundaries of said growth regions.
 14. Amethod as defined in claim 7, wherein said growth area of said substratehas a zero inclination relative to a basal plane.
 15. A method asdefined in claim 1, wherein said growth chamber is heated to atemperature of 800-2200° C.
 16. A method as defined in claim 1,including a step of using a non-reactive carrier gas to introduce atleast one of said crystal elements into said growth chamber.
 17. Amethod as defined in claim 1, wherein said crystal is an SIC crystal.18. A method of growing crystals containing a non-crystal element, of aconcentration of less than about 1.0×10¹⁶ cm⁻³, said crystal formed fromat least two crystal elements being grown by a process conducted in agrowth chamber, said crystal having at least two growth sites and saidnon-crystal element competitive for a growth site, said methodcomprising:a) identifying said growth site said non-crystal element iscompetitive for; b) identifying the crystal element which deposits insaid growth site said non-crystal element is competitive for; c)selecting a crystal element ratio of said two crystal elements whichcontains an amount of the competing crystal element which will competewith said non-crystal element during said crystal growth to form saidcrystal with said selected amount of non-crystal element; and d) flowingsaid selected amount of crystal elements into said growth chamber togrow said crystal.
 19. A method as defined in claim 18, wherein saidcrystal element ratio is selected to form a crystal having a non-crystalelement concentration of less than about 5.0×10¹⁵ cm⁻³.
 20. A method asdefined in claim 19, wherein said non-crystal element concentration ofless than about 1.0×10¹⁵ cm⁻³.
 21. A method as defined in claim 20,wherein said crystal element ratio is selected to form a crystal havinga non-crystal element concentration of less than about 8.0×10¹³ cm⁻³.22. A method as defined in claim 21, including a step of selecting asecond concentration of said non-crystal element; selecting a newcrystal element ratio to form said crystal with said secondconcentration of said non-crystal element; and changing said crystalelement ratio to said new crystal element ratio during the growing ofsaid crystal to form a crystal having multiple selected concentrationsof said non-crystal element.
 23. A method as defined in claim 18,wherein said crystal is an SIC crystal.
 24. A method containing anon-crystal element concentration of greater than about 1×10¹⁹ cm⁻³,said crystal formed from at least two crystal elements being grown by aprocess conducted in a growth chamber, said crystal having at least twogrowth sites and said non-crystal element competitive for a growth site,said method comprising:a) identifying said growth site said non-crystalelement is competitive for; b) identifying the crystal element whichdeposits in said growth site said non-crystal element is competitivefor; c) selecting a crystal element ratio of said two crystal elementswhich contains an amount of the competing crystal element which willcompete with said non-crystal element during said crystal growth to formsaid crystal with said selected amount of non-crystal element; and d)flowing said selected amount of crystal elements into said growthchamber to grow said crystal.
 25. A method as defined in claim 24,wherein said crystal element ratio is selected to form a crystal havinga non-crystal element concentration is greater than about 3.0×10¹⁹ cm⁻³.26. A method as defined in claim 25, wherein said crystal element ratiois selected to form a crystal having a non-crystal element concentrationis greater than about 5.0×10¹⁹ cm⁻³.
 27. A method as defined in claim26, wherein said crystal element ratio is selected to form a crystalhaving a non-crystal element concentration is greater than about1.0×10²⁰ cm⁻³.
 28. A method as defined in claim 27, wherein said crystalelement ratio is selected to form a crystal having a non-crystal elementconcentration is greater than about 5.0×10²⁰ cm⁻³.
 29. A method asdefined in claim 27, including a step of selecting a secondconcentration of said non-crystal element; selecting a new crystalelement ratio to form said crystal with said second concentration ofsaid non-crystal element; and changing said crystal element ratio tosaid new crystal element ratio during the growing of said crystal toform a crystal having multiple selected concentrations of saidnon-crystal element.
 30. A method as defined in claim 18, including astep of selecting a second concentration of said non-crystal element;selecting a new crystal element ratio to form said crystal with saidsecond concentration of said non-crystal element; and changing saidcrystal element ratio to said new crystal element ratio during thegrowing of said crystal to form a crystal having multiple selectedconcentrations of said non-crystal element.
 31. A method as defined inclaim 24, including a step of selecting a second concentration of saidnon-crystal element; selecting a new crystal element ratio to form saidcrystal with said second concentration of said non-crystal element; andchanging said crystal element ratio to said new crystal element ratioduring the growing of said crystal to form a crystal having multipleselected concentrations of said non-crystal element.
 32. A method asdefined in claim 24, wherein said crystal is an SIC crystal.
 33. Amethod of growing a crystal containing a selected ratio of at least onen-type dopant and at least one p-type dopant, said crystal formed fromat least two crystal elements being grown by a process conducted in agrowth chamber, said crystal having at least two growth sites and saidat least one n-type dopant primarily competitive for a n-type growthsite and said at least one p-type dopant primarily competitive for ap-type growth site, said method comprising of:a) identifying saidcrystal element which primarily deposits in said p-type growth site; b)identifying said crystal element which primarily deposits in said n-typegrowth site; c) selecting said ratio of n-type dopant to said p-typedopant incorporation; d) selecting an element ratio which contains acertain amount of said n-type growth site crystal element to said p-typegrowth site crystal element to obtain said amount of n-type dopant andp-type dopant to be deposited in said crystal during crystal growth,said ratio being dependent on said selected ratio of said n-type top-type dopant ratio; and e) flowing said at least two crystal elementsinto said growth chamber in said selected ratio amounts.
 34. A method asdefined in claim 33, wherein said crystal is an SIC crystal.
 35. Amethod of growing single crystals containing a selected concentration ofa non-crystal element, said single crystal formed from at least twocrystal elements being grown by a process conducted in a growth chamber,said single crystal having at least two growth sites and saidnon-crystal element competitive for a growth site, said methodcomprising:a) identifying said growth site said non-crystal element iscompetitive for; b) identifying the crystal element which deposits insaid growth site said non-crystal element is competitive for; c)controlling a crystal element ratio which contains a certain amount ofthe competing crystal element which will compete with said non-crystalelement during said crystal growth so that ≳1.0×10¹⁹ cm⁻³ of non-crystalelement is incorporated in said crystal.
 36. A method as defined inclaim 35, wherein said crystal is an SIC crystal.
 37. A method ofgrowing single crystals containing a selected concentration ofnon-crystal element, said single crystal formed from at least twocrystal elements being grown by a process conducted in a growth chamber,said single crystal having at least two growth sites and saidnon-crystal element competitive for a growth site, said methodcomprising:a) identifying said growth site said non-crystal element iscompetitive for; b) identifying the crystal growth which deposits insaid growth site said non-crystal element is competitive for; c)controlling a crystal element ratio which contains a certain amount ofthe competing crystal element which will compete with said non-crystalelement during said crystal growth so that ≲1.0×10¹⁵ cm⁻³ of non-crystalelement is incorporated in said single crystal.
 38. A method as definedin claim 37, wherein said crystal is an SIC crystal.
 39. A method ofselecting the amount of a dopant element deposited in a given growtharea of a single crystal and reducing the amount of a contaminantelement deposited in a given growth area of the single crystal, saidsingle crystal formed from at least two crystal elements as said crystalis grown at an area by a CVD process conducted in a growth chamber, saidcrystal elements comprising a first crystal element and a second crystalelement, said single crystal having at least two crystal growing siteswherein said first crystal element is deposited in a first growth siteand said second crystal element is deposited in a second element growthsite, said dopant element being competitive for at least one of saidfirst and second growth sites, said contaminant element beingcompetitive for a growth site other than said growth site of said dopantelement, said method comprising the steps of:a) identifying said growthsite said dopant element is competitive for; b) identifying said growthsite said contaminant element is competitive for; c) identifying a firstcrystal element which deposits in said growth site said dopant elementis competitive for; d) identifying a second crystal element whichdeposits in said growth side said contaminant element is competitivefor; e) selecting a ratio of said two crystal elements to obtain saidselected amount of said dopant element deposited in said competitivecrystal growth site and simultaneously reducing said amount of saidcontaminant element deposited in said competitive crystal growth site;and f) flowing said selected ratio of crystal elements into said growth.40. A method as in claim 39, including a step of changing the rate ofsaid dopant element deposited in said competitive growth site bychanging said ratio of said crystal elements being flowed into saidgrowth chamber.
 41. A method as defined in claim 40, wherein saidcrystal element ratio is selected to produce a crystal containing adopant element concentration of less than 1.0×10¹⁶ cm⁻³.
 42. A method asdefined in claim 41, wherein said dopant element concentration is lessthan 5.0×10¹⁵ cm⁻³.
 43. A method as defined in claim 40, wherein saidcrystal element ratio is selected to produce a crystal containing adopant element concentration of greater than 1.0×10.sup. cm⁻³.
 44. Amethod as defined in claim 43, wherein said dopant element concentrationis greater than 5.0×10¹⁹ cm⁻³.
 45. A method as defined in claim 40,including a step of varying at least once the ratio of said crystalelements being flowed into said growth chamber during the growing ofsaid crystal to form a crystal having at least two different dopantelement concentrations.
 46. A method as defined in claim 45, includingthe step of adding a second dopant element to reduce incorporation ofsaid contaminant element in said crystal, said second dopant elementhaving the same electrical polarity as said dopant element and saidsecond dopant competitive for said growth site said contaminant elementis competitive for.
 47. A method as defined in claim 46, including thestep of adding a second dopant element to reduce incorporation of saidcontaminant element in said crystal, said second dopant element havingthe same electrical polarity as said dopant element and said seconddopant competitive for said growth site said contaminant element iscompetitive for.
 48. A method as defined in claim 39, including a stepof varying at least once the ratio of said crystal element being flowedinto said growth chamber during the growing of said crystal to form acrystal having at least two different dopant element concentrations. 49.A method as defined in claim 48, including the step of adding a seconddopant element to reduce incorporation of said contaminant element insaid crystal, said second dopant element having the same electricalpolarity as said dopant element and said second dopant competitive forsaid growth site said contaminant element is competitive for.
 50. Amethod as defined in claim 39, including a step of adding a seconddopant element to reduce incorporation of said contaminant element insaid crystal, said second dopant element having the same electricalpolarity as said dopant element, and said second dopant competitive forsaid growth site said contaminant element is competitive for.
 51. Amethod as defined in claim 50, including a step of increasing the flowrate of said second dopant to further reduce said amount of saidcontaminant incorporated in said crystal.
 52. A method as defined inclaim 39, wherein said crystal is an SIC crystal.